Blends of oligopeptide terminal polyisobutylene or polystyrene

ABSTRACT

Various blends of polymers are disclosed, comprising oligopeptide functionalised polymers such as polyisobutylene and polystyrene. Mono-functionalised and di-functionalised polymers (each containing 0 to 5 peptide units beyond its terminal amide group) may be blended with each other and/or with non-functionalised polymers to produce blended compositions. Such compositions are of use, for example, in vibrations dampers. Certain blends also exhibit self-healing properties.

CROSS REFERENCE TO RELATED APPLICATIONS

This application is a continuation of co-pending U.S. application Ser.No. 14/646,950 which is a 371 National Stage of InternationalApplication No. PCT/EP2013/074793, filed Nov. 26, 2013, which claims thebenefit under 35 USC 119 of United Kingdom Application No. GB 1221246.0,filed Nov. 26, 2012.

FIELD OF THE INVENTION

This invention relates to polymer blends, in particular to blends ofpolymers with oligopeptide end groups (named oligopeptide-terminalpolymers in the following), and in particular to blends ofoligopeptide-terminal polyisobutylenes, polyisoprenes, polystyrenes andcopolymers of the aforementioned polymers. It also relates to the use ofmaterials from such blends, for example as reinforced, shape-persistentthermoplastic elastomers, as vibration damping materials, asself-healing materials and in other applications.

In typical thermoplastic materials, the polymer molecules start to flowalong each other upon heating, so that the material can bemacroscopically deformed. If this already happens at room temperature,albeit very slowly, the terms “creep” or “cold flow” are typicallydeployed. Materials exhibiting creep are not form-stable; they may beuseful in certain applications but inadequate for others whereshape-persistence is important.

In typical elastomers (rubbers), the polymer molecules are linkedtogether by chemical (covalent) bonds to form a three-dimensional,covalently linked network of chains with loops and holes. The chainsegments cannot move anymore, and on trying to deform the materialmacroscopically, it will retain its original shape upon release of thedeforming force (this is what makes the material a rubber). As covalentbonds only break when the material is destroyed at high temperatures,the covalent networks are irreversibly formed and stable.

High molecular weight polyisobutylene is widely used commercially as athermoplastic material. It shows elastomeric behaviour to a certaindegree due to chain entanglement but, unless covalently cross-linked(“butyl rubber”) shows a large degree of creep.

Interpenetrating polymer networks (IPNs) are elastomers that contain two(or more) independent, mechanically interlocked (interpenetrating)three-dimensional covalent networks consisting of different types ofpolymers. Firstly, one type of covalent network is formed from the firsttype of polymer chains by covalently connecting (cross-linking) them bya chemical reaction. Then the second type of polymer chains that fillthe loops and holes of the first network are cross-linked by a chemicalreaction. IPNs are useful as or in vibration damping materials. Withoutwishing to be bound by theory, we postulate that by subjecting such anIPN to a vibration of a certain frequency, only one of the twoindependent networks will start to respond (assuming the two types ofpolymers have different mechanical properties, e.g., stiffness) so thatat the molecular level, the segments of this network apparently start tomove along the segments of the other network, which dissipates energyand thus damps the vibration.

Mechanical vibration damping is highly desirable for structures, inorder to maintain their stability, performance and durability. Vibrationdamping is typically achieved by either “active damping”, using sensorsand actuators such as piezoelectric devices, or “passive damping” bymeans of materials that dissipate vibrational energy. Due to theirviscoelastic properties, polymer materials are useful passive dampingmaterials. Since the “loss factor” tan δ (defined as the ratio of lossand storage modulus, such as G″/G′ in shear mode) as a measure of thematerial's intrinsic damping properties attains its highest valuesaround the material's glass transition temperature, factors that controlthe latter are an important aspect of research in the field of dampingmaterials. For instance, small molecule additives have successfully beenused as plasticizers, in order to broaden the glass transition region.“Interpenetrating polymer networks” (IPNs), on the other hand, oftenfeature a superposition of the damping properties of the constitutinghomopolymers and, thus, offer opportunities for materials with excellentdamping performance over large temperature or frequency ranges. Finally,the incorporation of inorganic nanostructures or fibres as reinforcingfillers typically leads to a raising and broadening of the loss factormaximum around the glass transition temperature of the polymer matrix.This effect is particularly pronounced for fibres with a low aspectratio such as carbon whiskers or short microfibers. Moreover, it wasdemonstrated, that the interface properties between the matrix and thereinforcement strongly influence the energy dissipation mechanisms, withintermediate “interface friction” providing increased damping.

In contrast to true elastomers described above, “supramolecularmaterials” are an example of so-called “thermoplastic elastomers”(TPEs). In TPEs, the polymer chains are cross-linked (as perelastomers), but the network points (cross-links) are not covalent bonds(unlike elastomers). Instead, the network points are “weak” non-covalentbonds (secondary bonds, supramolecular interactions, such aselectrostatic interactions, hydrogen bonds, dipolar interactions, or Vander Waals interactions) that create stable network points but are weakenough to be reversibly broken/re-formed upon heating/cooling. Thismeans that the material is a three-dimensional network and behaves likean elastomer at temperatures below the critical temperature where thesupramolecular interactions start to be disrupted. Above thistemperature, the network is broken, the polymer chains can flow, and sothe material becomes thermoplastic and can be macroscopically deformed.This is advantageous because it enables processing the material into anydesired shape, and also from a sustainability standpoint allowing ameasure of re-use and recycling.

In some cases, the reversible disruption/re-formation of the networkscan even dynamically happen at room temperature (or under mild heatingor mechanical treatment). In this case, the materials may show a degreeof “self-healing”. This means that, by macroscopically damaging as bycutting for example, the material, the damage may heal just by itself,because molecules from both sides of the crack apparently can “leave”their network and reform a network with molecules from the other side.

In the prior art document WO 2011/045309 (BASF SE) individualoligopeptide-terminal polymers as fall within definition (a) herein and(b) herein (e.g. (a) monofunctional oligopeptide terminal polymers, and(b) difunctional oligopeptide terminal polymers) have been disclosedalong with their preparation and some postulated uses, but of anentirely different nature. However, in that prior patent, there is nodisclosure or suggestion of blends of different oligopeptide-terminatedpolymers or of blends of oligopeptide-terminated polymers with otherpolymers such as medium or high molecular weight PIB. Still less, isthere any suggestion of the particular advantageous properties,discovered now, attributable to such blends.

SUMMARY OF THE INVENTION

The invention provides, in one aspect, a polymer blend comprising: atleast two (a) monofunctional oligopeptide terminal polymers, or at leasttwo (b) difunctional oligopeptide terminal polymers.

The invention also provides a polymer blend comprising: at least onepolymer (a) as defined above, together with at least one polymer (b) asdefined above.

The invention further provides, in an alternative aspect, a polymerblend comprising: at least two (a) monofunctionaloligopeptide-terminated hydrophobic flexible polymers, or at least two(b) difunctional oligopeptide-terminated hydrophobic flexible polymers,or at least one of (a) together with at least one of (b) as definedabove.

In a further alternative aspect, the invention provides a polymer blendcomprising: at least two (a) monofunctional oligopeptide-terminatedhydrophobic polymers, or at least two (b) difunctionaloligopeptide-terminated hydrophobic polymers, or at least one of (a)together with at least one of (b) as defined above.

In some embodiments, the polymer blend may further comprise at least one(c) non-functionalised polymer, such as an isobutylene polymer. Thenon-functionalised polymer may have, for example, a molecular weightexceeding 10,000.

The invention also extends to a polymer blend comprising: at least onepolymer (c) as defined above, and at least one polymer (a) as definedabove and/or at least one polymer (b) as defined above.

In some embodiments, polymers (a), (b) and/or (c) comprise repeatingunits selected from isobutylene, butadiene, siloxane, acrylate,fluoropolymer, isoprene or styrene units. In some embodiments, thepolymer segment of polymer (a) and/or (b) and/or (c), may be selectedfrom the group consisting of: poly(isobutylene-co-isoprene),polyisoprene, polybutadiene, a □polysiloxane (in particularpoly(dimethylsiloxane)), a polyacrylate (in particular poly(methylacrylate) or poly(butyl acrylate)), poly(ethylene-co-butylene),hydrogenated poly(isoprene), hydrogenated poly(butadiene), or afluoropolymer (in particular poly(tetrafluoroethylene), andpoly(tetrafluoroethylene-co-ethylene).

Polymers (a) and (b) may, in some embodiments, comprise the same type ofpolymer segment. Where present, polymer (c) may, in some embodiments,comprise the same type of polymer segment as polymer (a) and/or (b).

In some particular embodiments, the blend comprisesoligopeptide-terminated isobutylene polymer with oligopeptide-terminatedstyrene polymer.

Polymers (a), (b) and/or (c) may comprise, for example, flexiblehydrophobic polymers.

In some embodiment, the oligopeptide-terminated polymer (a) may havefrom 0 to 5 (for example 2 to 5) peptide units beyond its terminal amidegroup and/or polymer (b) may have, at each end, 0 to 5 (for example 2 to5) peptide units beyond its terminal amide groups.

The oligopeptide moiety of polymer (a) and/or (b) may comprise, forexample, L-alanine units. In some embodiments, the oligopeptide moietyconsists only of such units.

The polymer segment of polymer (a) and/or (b) and/or, if present, (c),may, in some embodiments, be a hydrophobic polymer with a glasstransition temperature below 20° C.

Polymer blends according to the invention may, for example, be in theform of shape-persistent thermoplastic elastomers.

In some embodiments, the polymer blend may comprise interpenetratingsupramolecular polymer networks in which two or more specificsupramolecular interactions result in the formation of two or moreindependent, interpenetrating supramolecular networks with differenttransition (deaggregation) temperatures.

The invention also extends to vibration damping materials comprising thepolymer blends of the invention.

Some embodiments of vibration damping material comprise: (i)monofunctional oligopeptide terminal polyisobutylene, having from 2 to 5peptide units beyond its terminal amide group; and (ii) monofunctionaloligopeptide terminal polystyrene, having from 2 to 5 peptide unitsbeyond its terminal amide group.

The vibration damping material may, in some embodiments, additionallycomprise: (iii) a high molecular weight non-functionalised polymer.

Further particular embodiments of vibration damping material accordingto the invention include, by way of example—

-   -   A.—a material comprising: (i) monofunctional oligopeptide        terminal polyisobutylene, having from 2 to 5 peptide units        beyond its terminal amide group; and (ii) non-functionalised        polyisobutylene; wherein the monofunctional oligopeptide        terminal polyisobutylene may, for example, have 2 peptide units        beyond its terminal amide group; and    -   B.—a material comprising: (i) monofunctional oligopeptide        terminal polyisobutylene, having from 0 to 5 peptide units        beyond its terminal amide group; (ii) difunctional oligopeptide        terminal polyisobutylene, having from 0 to 5 peptide units        beyond its terminal amide group; and (iii) non-functionalised        polyisobutylene; wherein each of oligopeptide terminal        polyisobutylene polymers (i) and (ii) has the same number of        peptide units beyond its terminal amide group; wherein each of        oligopeptide terminal polyisobutylene polymers (i) and (ii) may,        for example, have 2 peptide units beyond its terminal amide        group.

Vibration damping materials according to the invention may, for example,be composite materials including one or more other components, such asone or more of the following: a plasticizer or a reinforcing filler(such as carbon fibre, carbon black, silica particles).

Vibration damping materials of the invention may be used, for example,in a form adapted to reduce vibration within a vehicle, such as in theform of a pad or other layer which can be interposed between members ofsuch vehicle subject to vibration. The invention also extends tovehicles which incorporate such damping materials, for example motorvehicles and aerospace vehicles.

In alternative aspects, the invention provides the use of any of thepolymer blends of the invention as a damping material and to a method ofvibration damping which involves the use of any of the polymer blends ofthe invention upon or within a structure such as a vehicle.

Moreover, the present invention also embraces polymer blends ofmonofunctional and/or difunctional oligopeptide-terminated polymersand/or the corresponding non-functional, higher molecular weightpolymers of flexible hydrophobic polymers other than polyisobutylene,such as polyisoprene, polybutadiene, polysiloxanes, polyacrylates, orfluoropolymers.

The present invention embraces, for example, blends of monofunctionaloligopeptide-terminated polyisobutylenes with oligopeptide segmentscomprising 0-5 amino acid repeating units (designated for convenienceM0-M5) and/or their difunctional analogues (similarly designated forconvenience as D0-D5) and/or oligopeptide-terminated polystyrene witholigopeptide segments comprising 0-5 amino acid repeating units(designated for convenience S0-S5) and/or oligopeptide-terminatedpolymers other than polyisobutylene/polystyrene (such as polybutadiene,polyisoprene, □hermogravime, polyacrylates, polymethacrylates, orcopolymers of any of the aforementioned polymers includingpolyisobutylene and polystyrene) with oligopeptide segments comprising0-5 amino acid repeating units (designated for convenience P0-P5) and/or(optionally high molecular weight) non-functionalised polyisobutylenesof different grades (molecular weights) and/or flexible hydrophobicpolymers other than polyisobutylene, such as polybutadiene,polyisoprene, polysiloxanes, polyacrylates, or fluoropolymers, orcopolymers of any of the aforementioned polymers includingpolyisobutylene and polystyrene.

The above blend(s), if required, can be used as a precursor blend to beadmixed with other non-functionalised, higher molecular weighthydrophobic flexible polymer such as polyisobutylene, to form thepreferred thermoplastic elastomer polymer blends of the invention.Alternatively, just one polymer (a) or (b) as defined herein may beadmixed with the non-functionalised, higher molecular weight hydrophobicflexible polymer, such as polyisobutylene.

These blends result in materials with novel properties because theoligopeptides deployed in the present invention form aggregates byhydrogen-bonding between the peptide groups; they are chiral; and their(exact) length determines what kind of aggregates form; and thislength-dependence is not only selective but even specific (self-sorting)so that different types of aggregates can persist in mixtures. Thepresent invention also embraces polymer blends using oligopeptideterminal groups based on other amino acids apart from those specificallyexemplified herein. Furthermore, whilst the present blends (materials)may be used on their own, they may also be used as the matrix materialfor a composite, that is, a reinforcing filler may be added into thepresent blends (such as carbon fibres, or carbon black, or silicaparticles). This may result in stronger, stiffer materials with the samebeneficial damping properties.

Further optional and preferred features are to be found amongst the subclaims herein:

(1) Some preferred embodiments of these blends are materials that are“Inherently Reinforced Thermoplastic Elastomers” in which nanostructuresformed by the aggregation of the oligopeptide segments act asreinforcing fillers. In particular, these materials are technologicallyadvantageous because they are thermoplastic materials with increasedmechanical moduli, low creep, and good thermal processability comparedto the corresponding high molecular weight polymers alone. See inparticular the manifold examples herein.

(2) Alternatively, further preferred embodiments of these blends resultin the formation of novel “Interpenetrating Supramolecular Networks”. Inparticular, these materials are technologically advantageous becausethey have excellent mechanical vibration damping properties. See inparticular the manifold examples herein.

(3) Still, further preferred embodiments of these blends can be deployedas materials with self-healing properties.

(4) Lastly, further embodiments of the invention are in the physicalform of composites which comprise the aforesaid polymer blends as wellas reinforcing fillers. Such composites can also be used as dampingmaterials with, for example, higher strength, stiffness, yet similardamping properties, showing a potential use of the present polymerblends materials as matrix materials for composites containing otherdesired or required components. These composites can also showself-healing properties.

The “Inherently Reinforced Thermoplastic Elastomer” embodiments can bederived from mixtures of molecules with oligopeptide termini of the sametype and length. These embodiments are examples of TPEs (in regard totheir properties) but are novel because they, due to the molecularlydefined oligopeptide end groups, already form networks at very shortsegment lengths of these end groups, so the inherent properties of theemployed polymer (polyisobutylene) do not alter too much. At the sametime, they are “reinforced” by the nanostructures (the tapes andfibrils) formed by the oligopeptides when they aggregate, similar toreinforcing a polymer with a filler (e.g., carbon fibres) to make ahigh-performance composite. As a result, we obtain materials withmechanical properties (moduli) matching or exceeding those of even highmolecular weight polyisobutylene, although we prefer to use very lowmolecular weight material. The present materials can be employed for thesame applications as high molecular weight polyisobutylene but can bethermally processed more easily, have better mechanical properties, showless creep, and are better from a sustainability standpoint (recyclingelastomers).

Some preferred embodiments of the aforesaid “Inherently ReinforcedThermoplastic Elastomer” embodiments can be derived from mixtures ofoligopeptide-terminated flexible polymers such as polyisobutylene witholigopeptide-terminated glassy polymers such as polystyrene or itscopolymers. In particular, these embodiments are novel because the phasesegregation of the immiscible polymers competes with the formation ofthe oligopeptide aggregates in the blends. Different from the microphasesegregation observed in typical block copolymers with domain sizes onthe order of tens of nanometers and above, this results in separateddomains of the immiscible flexible and amorphous polymers with diameterson the order of nanometers to tens of nanometers that are, in addition,connected by the oligopeptide aggregates. This creates a network ofnanoscopic oligopeptide aggregates and nanoscale glassy domains with aparticularly reinforcing effect, even at small weight fractions of theoligopeptide-terminated glassy polymers.

The “Interpenetrating Supramolecular Networks” embodiments are alsonovel and are obtained from mixtures of molecules with oligopeptidetermini of different type (oligopeptide sequence) or length, includingnon-functionalized, higher molecular weight polymers. In thesematerials, there are two independent networks (similar to IPNs), butboth are formed by cross-links that are non-covalent, weak, secondarybonds (e.g. hydrogen bonding). What this requires is that the networkformation relies on two “specific” (self-sorting) supramolecularinteractions that can form without interfering with each other. In thematerials provided by this invention, both types of networks rely on thesame type of supramolecular interactions (that is, hydrogen-bondingbetween the peptide functions) between two just differently longoligopeptides. The resulting networks can be either a supramolecular,hydrogen-bonded network of difunctional molecules (such as D0-D5); or a“percolation network” of tapes and fibrils formed byhydrogen-bond-driven aggregation of monofunctional molecules (M2-M5); ora percolation network of aforementioned tapes and fibrils with “hard”domains formed from oligopeptide-terminated glassy polymers (S2-S5 orother P2-P5); or an entanglement network formed from the high molecularweight polymers.

In any case, these embodiments share structural aspects of IPNs (twomechanically interlocked networks) and thermoplastic elastomers(reversibility of network formation). These embodiments thus extend uponregular supramolecular networks and other examples of TPE (that containno interpenetrating networks) and IPNs (that are formed from covalentnetworks), and as a result can function as excellent high performancedamping materials whilst simultaneously allowing beneficialthermoplastic processing (into different shapes). Their processing isalso flexible because it can be effected either above the temperaturewhere the first network melts, or above the temperature when also thesecond network melts, with different results. Specifically, it ispossible to heat above the melting temperature of both networks, processthe material into the desired shape (by injection moulding, extrusion,or other required technique), then cool below the melting temperature ofone network, let it form, then cool below the melting temperature of thesecond network to have it formed. In this way, an interpenetratingnetwork can be created just by processing, without using additionalchemical reaction steps (as is required for an IPN).

Some preferred embodiments of the aforesaid “InterpenetratingSupramolecular Networks” materials may also show self-healingproperties. One network may hold the material in place while the other(the weaker) network may dynamically break/re-form, either spontaneouslyor by heating it above that network's melting temperature, or by“mechanical treatment” (exposing the material to, e.g., a mechanicalvibration).

Advantageous properties of preferred embodiments of the present polymerblends include:

(1) Behaviour as reinforced thermoplastic elastomers, that is, rubberswith properties similar to or better than high molecular weightpolyisobutylene alone, but yet which can be processed thermoplasticallyby melting upon heating. They are ‘reinforced’ by the oligopeptidenanostructures, resulting in improved mechanical moduli (“strength”) andbehave as materials with no or low “creep” (cold flow; that is, theykeep their shape at room temperature). Commercial applications andindustrial uses can mirror those of regular PIB (e.g., barriermaterials) but their low creep and thermoplastic processing areunexpected advantages,

(2) Behaviour as materials with extremely large loss factors (losstangents tan delta) over large frequency and/or temperature ranges, i.e.apparent molecular level properties approaching liquids (at thosemechanical frequencies and temperatures) although they are in factsolids. This finds potential application in matrix materials forself-healing applications (materials that can cure mechanical damages‘themselves’),

(3) Significantly, excellent high-performance vibration dampingmaterials. Vibration damping uses being abundant in automotive andaerospace engineering whereby embodiments of the present blends findsubstantial application (even in their non-optimized state, someembodiments of the present invention already match or out-performoptimised multi-component commercial formulations.

The materials disclosed in the present invention implement usefulvibration damping properties in a novel way. Due to the formation of theaforesaid “interpenetrating supramolecular networks”, these materialsthemselves combine the beneficial effects (with respect to energydissipation upon mechanical excitation as needed for vibration damping)of low aspect ratio reinforcing fillers, interpenetrating networks, andlow molecular weight poly(isobutylene) plasticizers. They, therefore,yield shape-persistent materials with excellent energy dissipation anddamping properties over large frequency and temperature ranges, withoutthe use of additional components such as additional fillers orplasticisers.

BRIEF DESCRIPTION OF THE DRAWINGS

In order that the invention may be further described, more easilyappreciated and readily carried into effect by those skilled in the art,reference will now be made to embodiments by way of non-limiting exampleonly and with reference to the accompanying drawings, in which:

FIGS. 1A, 1B and 1C represent schematic illustrations of the selectiveself-assembly of the monofunctional oligo(L-alanine)-modifiedpoly(isobutylene)s M0-M5 (n=0-5; x≈20) and the correspondingdifunctional derivatives D0-D5 (n=0-5; x 20), FIG. 1D represents aschematic illustration in which the coexistence of these nanostructuresin blends of molecules with matching oligopeptide termini results in‘inherently reinforced thermoplastic elastomers’, and FIG. 1E representsa schematic illustration in which the coexistence of thesenanostructures in blends of molecules with different oligopeptidetermini results in ‘interpenetrating supramolecular networks’,

FIGS. 2A, 2B, 2C and 2D represent amide A and amide I regions of thesolid-state infrared (IR) spectra of bulk samples of PIB-Ala_(n)-AcM0-M5 as well as Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5, and FIGS. 2E and 2Frepresent corresponding plots of the position of the global maxima ofthe amide A and amide I absorptions,

FIGS. 3A, 3B, 3C and 3D represent amide A and amide I regions of thesolution-phase infrared (IR) spectra of samples of PIB-Ala_(n)-Ac M0-M5as well as Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 in dilute solution intetrachlorethane, and FIGS. 3E and 3F represent corresponding plots ofthe position of the global maxima of the amide A and amide Iabsorptions,

FIG. 4A represents peak deconvolution of the amide I regions of samplesof PIB-Ala_(n)-Ac M0-M5 in bulk, FIG. 4B represents peak deconvolutionof the amide I regions of samples of Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 inbulk, FIG. 4C represents peak deconvolution of the amide I regions ofsamples of PIB-Ala_(n)-Ac M0-M5 in dilute solution in tetrachlorethane,and FIG. 4D represents peak deconvolution of the amide I regions ofsamples of Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 in dilute solution intetrachlorethane,

FIG. 5 shows atomic force microscopy (AFM) images of M1-M5 spin-coatedfrom tetrachlorethane solution onto either SiO₂ or HOPG substrates,

FIGS. 6A and 6B show thermogravimetric analysis of M0-M5 and D0-D5 aswell as the parent poly(isobutylene)s PIB-NH₂ and H₂N—PIB-NH₂, FIG. 6Crepresents differential scanning calorimetry of M0-M5 as well as theparent poly(isobutylene) PIB-NH2, and FIG. 6D represents differentialscanning calorimetry of D0-D5 as well as the parent poly(isobutylene)H2N-PIB-NH2,

FIGS. 7A and 7B represent amide I regions of the temperature-dependentsolid-state IR spectra of M2 and D2,

FIGS. 8A, 8B and 8C show rheological dynamic frequency sweep experimentsat 25° C. of unmodified PIB of different molecular weights (1200 forPIB-NH₂, 2500 for H₂N-PIB-NH₂, 35'000, 75'000, 200'000, 425'000),showing a) storage moduli G′, b) loss moduli G″, and c) viscosity 2 andD2,

FIGS. 9A, 9B, 9C, 9D, 9E and 9F show rheological dynamic frequency sweepexperiments at 25° C. of PIB-Ala_(n)-Ac M0-M5 andAc-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 as well as the parent poly(isobutylene)s1 and 4, showing a,d) storage moduli G′, b,e) loss moduli G″, and c,f)viscosity,

FIGS. 10A, 10B and 10C show rheological dynamic frequency sweepexperiments at 25° C. of different binary blends M2/D2, showing a)storage moduli G′, b) loss moduli G″, and c) and viscosity,

FIGS. 11A, 11B, 11C and 11D shows a comparison of storage (G′) and lossmoduli (G″) at 1 rad/s of a) unmodified PIB as a function of molecularweight, c,d) M0-M5 and D0-D2 as a function of hydrogen-bonding sites perend group, and b) of different blends of M2/D2 as a function ofcomposition,

FIGS. 12A and 12B represents amide A, amide I and amide II regions ofthe solid-state infrared (IR) spectra of M2, D2, as well as the completeseries of M2/D2 blends (Examples 1-6),

FIGS. 13A and 13B represent amide I regions of the temperature-dependentsolid-state IR spectra of the blend M2/D2 9:1 (Example 3), and FIG. 13Cshows the parent PIB amine was a viscous liquid, M2 was a sticky solid,D2 a brittle powder, the blend M2/D2 9:1 (Example 3) was an ‘inherentlyreinforced’ thermoplastic elastomer,

FIGS. 14A, 14B and 14C show temperature-dependent shear rheology of theblend M2/D2 9:1 (Example 3) and FIGS. 14D, 14E and 14F showtemperature-dependent shear rheology of PIB (MW 200'000) as a referencematerial,

FIGS. 15A, 15B and 15C show Temperature-dependent shear rheology of theblend M2/D1 1:4 (Example 7), compared to pure M2, pure D1, as well asthe blend M2/D2 9:1 (Example 3) including plots of a) storage modulusG′, b) loss modulus G″, and c) viscosity, and FIG. 15D representsdifferential scanning calorimetry of M2/D1 1:4,

FIGS. 16A, 16B and 16C show rheological time-temperature superposition(TTS) master curves of D1, M2, and M2/D1 1:4 (Example 7),

FIGS. 17A and 17B show amide A, amide I and amide II regions of thesolid-state infrared (IR) spectra of solution phase IR spectra ofsamples in tetrachlorethane and FIGS. 17C and 17D show amide A, amide Iand amide II regions of the solid-state infrared (IR) spectra of solidstate IR spectra of bulk samples of PIB (MW 75'000), PIB (MW 35'000),M2, D2, as well as different binary and ternary blends of M2 and D2 inPIB (MW 75'000) or PIB (MW 35'000) (Examples 8-10),

FIG. 18A represents differential scanning calorimetry of M2, D2 andtheir binary and ternary blends M2/PIB (MW 35'000) 5:5 (Example 9) andM2/D2/PIB (Mw 35'000) 4:1:5 (Example 10), and FIG. 18B showsthermogravimetric analysis of M2, D2 and their binary and ternary blendsM2/PIB (MW 35'000) 5:5 (Example 9) and M2/D2/PIB (Mw 35'000) 4:1:5(Example 10) as well as for PIB-NH2 and H2N-PIB-NH2,

FIG. 19 shows atomic force microscopy (AFM) height (left) and phase(right) images of M2 and M2/PIB (MW 35'000) 5:5 (Example 9),

FIGS. 20A, 20B, 20C, 20D, 20E and 20F show rheological time-temperaturesuperposition (TTS) master curves of M2/PIB (MW 75'000) 5:5 (Example 8),M2/PIB (MW 35'000) 5:5 (Example 9), M2/D2/PIB (MW 35'000) 4:1:5 (Example10),

FIGS. 21A, 21B and 21C show an experimental setup for a random vibrationmodal analysis test on a sandwich structure comprising a damping layer,and

FIGS. 22A, 22B, 22C, 22D and 22E represent Lissajous curves obtainedfrom oscillatory shear stress-strain test, performed on a rheometer, andFIGS. 22F and 22H represent plots of the logarithm of the dissipatedenergies relative to the logarithm of the strain applied at −45° C.during an oscillatory shear stress-strain test, performed on arheometer.

FIGS. 23A and 23B show atomic force microscopy (AFM) height images anddifferential scanning calorimetry measurements of the blend of M2 withunmodified polystyrene (9:1) and the blend M2/S2 9:1 (Example 11); FIG.23C shows atomic force microscopy (AFM) height images and FIG. 23D showsdifferential scanning calorimetry measurements of the blend M3 withunmodified polystyrene (9:1) or the blend M3/S3 9:1 (Example 12).

FIGS. 24A and 24B show storage moduli G′ and loss moduli G″ determinedby rheological dynamic frequency sweep experiments at 25° C. of M2, S2,the blend of M2 with unmodified polystyrene (9:1) and the blend M2/S29:1; FIGS. 24C and 24D show storage moduli G′ and loss moduli G″determined by rheological dynamic frequency sweep experiments at 25° C.of M3, S3, the blend of M3 with unmodified polystyrene (9:1) and theblend M3/S3 9:1.

FIG. 25 shows a rheological time-temperature superposition master curveof the ternary blend M3/S3/PIB (MW 35'000) 9:3:12 (Example 13).

DETAILED DESCRIPTION OF EXAMPLES

As shown in the drawings and referring in particular to FIGS. 1A, 1B, 1Cand 1D:

FIGS. 1A, 1B and 1C provide Schematic illustration of the selectiveself-assembly of the monofunctional oligo(L-alanine)-modifiedpoly(isobutylene)s M0-M5 (n=0-5; x 20) and the correspondingdifunctional derivatives D0-D5 (n=0-5; x 20) into small hydrogen-bondedaggregates, flexible single β-sheet tapes, or rigid stacked β-sheetfibrils. The coexistence of these nanostructures in blends of moleculeswith different oligopeptide termini resulted in ‘inherently reinforcedthermoplastic elastomers’ (see FIG. 1D) or ‘interpenetratingsupramolecular networks’ (see FIG. 1E).

Differently from all previous examples of supramolecular networks, theaggregation of the oligopeptide-terminated polymers that constitute thebasis of the present invention that comprise chiral and monodisperse(molecularly defined) oligopeptides as hydrogen-bonded ligands resultsin a highly selective formation of small hydrogen-bonded aggregates fromcompounds with short oligopeptides (such as M0-M1, D0-D1), flexiblesingle β-sheet tapes from compounds with medium-size oligopeptides (suchas M2-M3, D2-D3), or rigid stacked β-sheet fibrils from compounds withlonger oligopeptides (such as M4-M5, D4-D5), because the helicalconformation of single oligopeptide β-strands, the induced helicaltwisting of β-sheets, and finally the number of stacked β-sheets areintimately interrelated. This length-dependent self-assembly is even“self-sorting”, that is, specific in the sense that the differentnanostructures obtained from different oligopeptide segments coexist inbulk. It is this particular feature that has enabled us to tailor thethermomechanical properties of the blends. Thus, mixtures of moleculeswith “matching” oligopeptide termini (identical oligopeptide length andamino acid sequence) gave rise to thermoplastic elastomers that were“inherently reinforced” with β-sheet tapes or fibrils. By contrast,blends of derivatives with “non-matching” oligopeptide termini(different oligopeptide length or amino acid sequence, includingnon-functionalized polymers) formed novel “interpenetratingsupramolecular networks”. It is worth noting that in both cases, networkformation allows for dynamic network reorganization processes and maygive rise to self-healing or thermoresponsive materials. In this regard,polyisobutylene soft segments have proven to be of high interest, due totheir conformational dynamics and resulting macroscopic properties. Seein particular the manifold examples. Further examples of blends ofnon-functionalised and oligopeptide-terminated derivatives of flexibleand hydrophobic polymers, such as polyisoprene, polybutadiene,polyacrylates, polysiloxanes, or fluoropolymers, share the samestructural features and properties and are embraced by the currentinvention.

Referring in particular to FIGS. 2A-5:

FIGS. 2A, 2B, 2C and 2D provide Amide A and amide I regions of thesolid-state infrared (IR) spectra of bulk samples of PIB-Ala_(n)-AcM0-M5 as well as Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5, as well as FIGS. 2Eand 2F provide corresponding plots of the position of the global maximaof the amide A and amide I absorptions as a function of the number ofalanine repeating units n revealed that M2-M5 and D2-D5 exhibited asingle amide A absorption at 3270-3276 cm⁻¹, a strong and sharp amide Iabsorption at 1625-1627 cm⁻¹ (half-height width≈16-17 cm⁻¹), and a sharpsecondary absorption at 1687-1695 cm⁻¹, all consistent with the presenceof highly ordered antiparallel β-sheet structures. See FIGS. 4A and 4Bfor peak deconvolutions.

FIGS. 3A, 3B, 3C and 3D provide Amide A and amide I regions of thesolution-phase infrared (IR) spectra of samples of PIB-Ala_(n)-Ac M0-M5as well as Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 in dilute solution intetrachlore thane, as well as FIGS. 3E and 3F provide correspondingplots of the position of the global maxima of the amide A and amide Iabsorptions as a function of the number of alanine repeating units nrevealed that, in solution, M0-M2 and D0-D2 remained non-aggregated. Asharp transition was then observed for longer oligopeptides; M3-M5 aswell as D3-D5 exhibited a single amide A absorption at 3271-3273 cm−1, astrong and sharp amide I absorption at 1624-1625 cm−1 (half-heightwidth≈14−17 cm−1), and a sharp secondary absorption at 1690-1695 cm−1,all consistent with the presence of highly ordered antiparallel β-sheetstructures in solution. See FIGS. 4C and 4D for peak deconvolutions.

Peak deconvolution of the amide I regions of samples of PIB-Ala_(n)-AcM0-M5 as well as Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 are shown in bulk (seeFIGS. 4A and 4B) and in dilute solution in tetrachlorethane (see FIGS.4C and 4D); global maxima labelled in blue; β-sheet bands in red;predominant bands in black. Although the peak fitting was started withthe same number of bands at approximately the same positions, limitingtheir width to reasonable values in all cases, the results of thedeconvolution were still sensitive to the exact starting parameters and,therefore, just served to obtain an estimate for the peak area A1625 ofthe absorption bands at around 1625-1630 cm⁻¹ relative to the total peakarea AI, total of the amide I absorption. The latter is a qualitativeassessment for the relative degree of aggregation. See FIG. 2d (maintext) for a plot of A1625/AI, total as a function of the number ofalanine repeating units n.

FIG. 5. Atomic force microscopy (AFM) images of M1-M5 spin-coated fromtetrachlorethane solution onto either SiO₂ or HOPG substrates revealedthe formation of fibrils for M5; mixtures of fibrils and tapes for M4 onHOPG and mixtures of fibrils and drop-like features on SiO₂; no definedaggregates for M0-M3 on SiO₂, but long tapes for M3, short laterallyaggregated tapes for M2, and continuous films for M1 and M0 on HOPG.

The monofunctional compounds M0-M5 and the difunctional compounds D0-D5exhibited distinctly length-dependent aggregation properties. Accordingto IR spectroscopy, M4-M5 and D4-D5 gave rise to highly ordered andstrongly aggregated antiparallel β-sheet structures both in bulk and insolution. M2-M3 and D2-D3 were only aggregated in bulk materials. Theend groups in M0-M1 and D0-D1 were too short to induce β-sheet formationeither in the bulk or in solution. Atomic force microscopy (AFM) imagingthen established a link to the corresponding nanoscopic morphologies forthe monofunctional derivatives M0-M5. Thus, rigid and many micrometreslong fibrils with diameters of a few nanometres were observed for M5 andM4 on both highly oriented pyrolytic graphite (HOPG) and SiO₂substrates. The dimensions of fibrils obtained from M5 suggested thatthey were formed from 4-6 stacked β-sheet tapes. In the case of M4, thefibrils were formed from 2-4 stacked β-sheet tapes, according to theircross-sections determined by AFM imaging. M3 gave rise to long flexiblefibrils or tapes on HOPG that were thinner than those of M4 andexhibited an epitaxial orientation with the substrate. In the case ofM2, we observed laterally aggregated tape-like features on HOPG withlengths on the order of a few hundred nanometres. The epitaxialorientation of the tape-like features from M2-M4 on HOPG as well astheir absence on SiO₂ substrates suggested that they had not alreadybeen present in solution but formed upon drying of the sample on the AFMsubstrate, in agreement with the IR spectroscopic results. Hence, ourresults prove that longer oligopeptides did not only result in theexpected increase in aggregation strength but that superstructureformation was also affected, due to the molecular chirality of anoligo(L-alanine) segment. We proved that we selectively obtained rigidstacked β-sheet fibrils from the “longest” oligopeptides (n 4 alanineresidues); single β-sheet tapes from “medium-size” oligopeptides (n=2-3alanine residues) in the bulk; and weak, undefined aggregates from shorthydrogen-bonded end groups (n=0-1 alanine residues).

Referring in particular to FIGS. 6A-7B:

FIGS. 6A and 6B providing Thermogravimetric analysis of M0-M5 and D0-D5as well as the parent poly(isobutylene)s PIB-NH₂ and H₂N—PIB-NH₂revealed that complete PIB depolymerization occurred at temperaturesabove 340° C. in all cases. Whereas derivatives M0-M2 and D0-D2 werestable up to temperatures of at least 250° C., compounds with longeroligopeptides noticeably underwent a first stage of decompositionalready at temperatures of around 170-200° C., tentatively assigned to adegradation of the oligopeptides, supposedly by ring-closingfragmentation. FIGS. 6C and 6D providing Differential scanningcalorimetry revealed that only M2 and D2 exhibited reversible thermaltransitions at 170° C. (15.5 J/g) and 178° C. (16 J/g), respectively.These could be assigned to the ‘melting’ (deaggregation) and‘crystallization’ (aggregation) of the β-sheet aggregates (see FIGS. 7Aand 7B). Derivatives with longer oligopeptides showed endothermic peaksat increasingly high temperatures that were already in the range of orabove their decomposition temperatures and, accordingly, did not exhibitany exothermic peaks upon cooling, except for M3 which exhibited a weakexothermic peak at 238° C. (3 J/g).

FIGS. 7A and 7B providing Amide I regions of the temperature-dependentsolid-state IR spectra of M2 and D2 revealed that the position andintensity of the absorption bands associated with β-sheet secondarystructures remained virtually unchanged until at least 150° C. but thenrapidly decreased at temperatures above 170° C., proving that thereversible transitions observed in DSC (see FIGS. 6A, 6B, 6C and 6D)were associated to ‘melting’ (deaggregation) and ‘crystallization’(aggregation) of the β-sheet aggregates.

Thermogravimetric analysis (TGA), differential scanning calorimetry(DSC), and temperature-dependent solid-state IR spectroscopy proved thatM0-M2 and D0-D2 were straightforwardly processable below theirdegradation temperature of 250° C. Moreover, M2 and D2 exhibiteddetectable reversible thermal transitions at 170° C. and 178° C.,respectively, according to DSC, that solid state IR spectroscopy provedto be associated with β-sheet deaggregation.

Referring in particular to FIGS. 8-9:

FIGS. 8A, 8B and 8C provide Rheological dynamic frequency sweepexperiments at 25° C. of unmodified PIB of different molecular weights(1200 for PIB-NH₂, 2500 for H₂N—PIB-NH₂, 35'000, 75'000, 200'000,425'000), showing storage moduli G′ (FIG. 8A), loss moduli G″ (FIG. 8B),and viscosity |η*| (FIG. 8C); the shear viscosity formally defined as|η*|=p_(21,0)/(γ₀·ω) was used as a calculated entity to compare bothliquid and rubbery materials. The γ (% strain) values used for theseexperiments ranged from 3-30% and were selected to be in the linearregime of the investigated materials. Depending on their molecularweight, the materials undergo a transition in mechanical properties,from low viscosity liquid behaviour with a zero-shear viscosity of 32 Pas for PIB-NH₂ and 140 Pa s for H₂N—PIB-NH₂ to, finally, rubberybehaviour for PIB (MW 200'000 and 425'000). Hence, the latter twomaterials exhibited a frequency-dependent viscosity |η*| with a slope of−1, and both their storage (G′) and loss moduli (G″) at 1 rad/s (as areference) were increased by 7 and 3 orders of magnitude, respectively,as compared to low molecular weight PIB-NH₂.

FIGS. 9A-9F providing Rheological dynamic frequency sweep experiments at25° C. of PIB-Ala_(n)-Ac M0-M5 and Ac-Ala_(n)-PIB-Ala_(n)-Ac D0-D5 aswell as the parent poly(isobutylene)s 1 and 4, showing storage moduli G′(FIGS. 9A and 9D), loss moduli G″ (FIGS. 9B and 9E), and viscosity |η*|(FIGS. 9C and 9F); the shear viscosity formally defined as|η*|=p_(21,0)/(γ₀·ω) was used as a calculated entity to compare bothliquid and rubbery materials. The γ (% strain) values used for theseexperiments ranged from 0.03-30% and were selected to be in the linearregime of the investigated materials. Depending on the number ofhydrogen-bonding sites in the series M0-M5, the materials showed atransition from moderately viscous liquid behaviour to a rubberybehaviour for M3-M5. The latter materials exhibited afrequency-dependent viscosity |η*| with a slope of −1 and both theirstorage (G′) and loss moduli (G″) at 1 rad/s (as a reference) wereincreased by 7 and 3 orders of magnitude, respectively, as compared tothe constituent low molecular weight PIB-NH₂ 1. Difunctional derivativesD0-D2 can give rise to supramolecular networks and, hence, showed aneven more drastic transition of mechanical properties. Thus, D2 wasalready brittle hard solid and exhibited a storage modulus of G′=2 MPa,10 times higher than high molecular weight poly(isobutylene)(MW≥200'000). The pure higher difunctional homologues D3-D5 were hardand brittle powders that could not be processed into solid discs and,hence, were not investigated by means of shear rheology.

Shear rheology on the monofunctional derivatives M0-M5 in comparison tounmodified polyisobutylenes revealed a transition of mechanicalproperties from moderately viscous liquid (M0-M1) to a rubbery behaviour(M4-M5) as a function of the number of n alanine residues. Starting withthe β-sheet tape-forming derivatives M2-M3, the materials exhibitedshear moduli exceeding those of high molecular weight polyisobutylenes,although the molecular weight of the attached polymer (MW 1,200) was farbelow the entanglement length of PIB (≈15,000) and the monofunctionalderivatives cannot form hydrogen-bonded networks. The storage and lossmoduli at 1 rad/s within the series leveled off towards G′≈0.6 MPa andG″≈0.06 MPa, indicating that a further increase of the oligopeptidelength would not substantially affect the materials' mechanicalproperties anymore. The network-forming difunctional derivatives D0-D2showed even more drastic changes in mechanical properties as a functionof oligopeptide length, and D2 was already a brittle hard solid (G′=2MPa). While the notion of a mechanical reinforcement is certainly wellin line with previous examples of supramolecular networks, the observeddrastic dependence of aggregation behaviour and mechanical properties onthe number of hydrogen-bonding sites allows for further tailoring of thematerials' thermomechanical properties in blends of the investigatedcompounds. See in particular the disclosed examples.

Referring in particular to FIGS. 10A-14E (Examples 1-6):

FIGS. 10A, 10B and 10C provide Rheological dynamic frequency sweepexperiments at 25° C. of different binary blends M2/D2, showing a)storage moduli G′ (FIG. 10A), b) loss moduli G″ (FIG. 10B), and c) andviscosity |η*| (FIG. 10C); the shear viscosity formally defined as|η*|=p_(21,0)/(γ₀·ω) was used as a calculated entity to compare bothliquid and rubbery materials. The γ (% strain) values used for theseexperiments ranged from 0.05-0.5% and were selected to be in the linearregime of the investigated materials. Large changes in the moduli andviscosities were observed between pure M2 and blends of up to 10 wt % ofD2 in M2, which then level off for compounds with higher content of D2.

FIGS. 11A, 11B, 11C and 11D provide a Comparison of storage (G′) andloss moduli (G″) at 1 rad/s of unmodified PIB as a function of molecularweight (FIG. 11A), M0-M5 and D0-D2 as a function of hydrogen-bondingsites per end group (FIGS. 11C and 11D), and of different blends ofM2/D2 as a function of composition (FIG. 11B). Storage moduli of M3-M5and D2 as well as blends of at least 5% D2 in M2 (G′=0.6 MPa) exceedthose of high molecular weight PIB. Blends M2/D2 with more than 10 wt %of D2 show storage and loss moduli levelling off toward G′≈2 MPa andG″≈0.1 MPa as observed for pure D2. In the series of M0-M5, the storageand loss moduli appear to converge toward G′≈0.6 MPa and G″≈0.06 MPa.

FIGS. 12A and 12B provide Amide A, amide I and amide II regions of thesolid-state infrared (IR) spectra of M2, D2, as well as the completeseries of M2/D2 blends (Examples 1-6). They all exhibited a single amideA absorption at 3275 cm⁻¹, a strong and sharp amide I absorption at 1627cm⁻¹, a smaller absorption at 1687 cm⁻¹ and an amide II absorption at1543 cm⁻¹, consistent with the presence of highly ordered antiparallelβ-sheet structures. Moreover, independent of their composition, theblends exhibited and amide I regions indistinguishable from the purecompounds, providing evidence for the presence of antiparallel β-sheetstructures and the miscibility of D2 in M2.

FIGS. 13A and 13B providing Amide I regions of the temperature-dependentsolid-state IR spectra of the blend M2/D2 9:1 (Example 3) revealed thatthe position and intensity of the absorption bands associated withβ-sheet secondary structures remained virtually unchanged until at least150° C. but then rapidly decreased at temperatures above 170° C.,exactly like the individual components M2 and D2. This proves that thethermal transitions were associated to ‘melting’ (deaggregation) and‘crystallization’ (aggregation) of the β-sheet aggregates and that thetwo compounds formed a common hydrogen-bonded network together. FIG. 13Cshows that whereas the parent PIB amine was a viscous liquid, M2 was asticky solid, D2 a brittle powder, the blend M2/D2 9:1 (Example 3) wasan ‘inherently reinforced’ thermoplastic elastomer.

FIGS. 14A, 14B, 14C, 14D, 14E and 14FE show Temperature-dependent shearrheology of the blend M2/D2 9:1 (Example 3) (FIGS. 14A, 14B and 14C) andPIB (MW 200'000) as a reference material (FIGS. 14D, 14E and 14F). Plotsof storage modulus G′, loss modulus G″, and viscosity |η*| (formallydefined as |η*|=p_(21,0)/(γ₀·ω) and used as a calculated entity tocompare both liquid and rubbery materials) at 1 rad/s as a function oftemperature and examples of rheological dynamic frequency sweepexperiments at 25° C. and 180° C. showed that M2/D2 9:1 (Example 3)experienced a sharp and single-step decrease of its moduli and viscosityat a temperature of about 160° C., yielding low viscosity liquids, thathad been found to coincide with β-sheet deaggregation according totemperature-dependent IR spectroscopy (see FIG. 13a ). By contrast, PIB(MW 200'000) remained in the rubbery state up to temperatures of atleast 250° C. The γ (% strain) values used for these experiments rangedfrom 0.1-50% and were selected to be in the linear regime of theinvestigated materials.

Binary blends of monofunctional and difunctional derivatives with“matching” oligopeptide segments were found to give rise tosupramolecular networks that were “inherently reinforced” by theincorporated β-sheet aggregates. Specifically, binary blends of thethermally processable compounds M2 and D2 with the compositions (byweight) M2/D2 99:1 (Example 1), 95:5 (Example 2), 9:1 (Example 3), 7:3(Example 4), 5:5 (Example 5), and 1:9 (Example 6) were obtained bydissolving mixtures of the compounds in tetrachlorethane (TCE), stirringthe solutions at room temperature for 16 h, removing the solvent invacuo, and drying the resulting materials in high vacuum at 120° C. for3 days. Independent of their composition, the blends exhibited solidstate IR spectra with amide I regions indistinguishable from the purecompounds and underwent a single-step “melting” transition at 160-170°C. associated with the deaggregation of all β-sheet structures. Theblends yielded rubbery materials with shear moduli that exceeded thoseof even high molecular weight PIB (0.2 MPa) by an order of magnitudeeven for low fractions of D2. Thus, the storage moduli already reachedG′=0.6 MPa upon the addition of ≥5 wt % D2 (Examples 2-6) and leveledoff toward G′ 2 MPa for 10 wt % D2 (Examples 3-6). At the same time, thelatter materials (Examples 3-6) experienced a sharp decrease of theirmoduli and viscosities at their melting temperatures. Hence, we obtained“inherently reinforced” polyisobutylene-based thermoplastic elastomersthat exhibited superior shear properties and showed lower creepbehaviour at room temperature, but yielded well-processable melts atelevated temperatures, well below their decomposition temperature.

Table 1 shows representative values of storage moduli G′, loss moduliG″, loss factors tan δ, and viscosities |η*| for different grades ofpolyisobutylenes, M0-M5, D0-D2, as well as Examples 1-6.

Referring in particular to FIGS. 15A-22H (Examples 7-10):

FIGS. 15A, 15B and 15C show Temperature-dependent shear rheology of theblend M2/D1 1:4 (Example 7), compared to pure M2, pure D1, as well asthe blend M2/D2 9:1 (Example 3). Plots of a) storage modulus G′ (FIG.15A), b) loss modulus G″ (FIG. 15B), and c) viscosity |η*| (formallydefined as |η*|=p_(21,0)/(γ₀·ω) was used as a calculated entity tocompare both liquid and rubbery materials) (FIG. 15C) at 1 rad/s as afunction of temperature revealed that the material underwent a two-stagethermomechanical transition, first following the behaviour of D1 in thetemperature range of −45° C. to above room temperature, and then M2between 65° C. and the melting transition at 139° C. FIG. 15D providingDifferential scanning calorimetry showed that the blend M2/D1 9:1(Example 7) exhibited a transition at about 25° C. (assigned to themelting of the D1 network) as well as a reversible transitions at 139°C. (onset at 128° C.; 0.5 J/g) that we assigned to the reversibledeaggregation of the β-sheet aggregates of M2. While the apparent‘melting point depression’ as compared to pure M2 (170° C.) suggests acertain interaction between D1 and M2 in that temperature range, thepronounced effect of the minority component M2, the two-stagetemperature transition, and the superimposed rheological properties ofthe pure components in the blend provide sufficient evidence for thepresence of two independent, ‘interpenetrating supramolecular networks’that do not undergo macrophase segregation.

FIGS. 16A, 16B and 16C provide Rheological time-temperaturesuperposition (TTS) master curves of D1, M2, and M2/D1 1:4 (Example 7)at T_(ref)=25° C. D1 showed an entanglement point at 25° C. M2 exhibiteda large tan δ peak at an unusual temperature as compared to highmolecular weight PIB. M2/D1 1:4 (Example 7) possessed a broad regionwith pronounced ‘liquid-like’ behaviour and a large loss factor of up totan δ=2.0.

FIGS. 17A, 17B, 17C and 17D provide Amide A, amide I and amide IIregions of the solid-state infrared (IR) spectra of solution phase IRspectra of samples in tetrachlorethane (FIGS. 17A and 17B) and soldstate IR spectra of bulk samples of PIB (MW 75'000), PIB (MW 35'000),M2, D2, as well as different binary and ternary blends of M2 and D2 inPIB (MW 75'000) or PIB (MW 35'000) (Examples 8-10) (FIGS. 17C and 17D).All mixtures were deaggregated in solution, but bulk materials exhibiteda single amide A absorption at 3276 cm⁻¹, a strong and sharp amide Iabsorption at 1627 cm⁻¹, a smaller absorption at 1687 cm⁻¹ and an amideII absorption at 1543 cm⁻¹, consistent with the presence of highlyordered antiparallel β-sheet structures. Moreover, independent of theircomposition, the blends exhibited and amide I regions indistinguishablefrom the pure compounds, providing evidence for the presence ofantiparallel β-sheet structures dispersed in a PIB matrix.

FIG. 18A providing Differential scanning calorimetry revealed that M2,D2 and their binary and ternary blends M2/PIB (MW 35'000) 5:5 (Example9) and M2/D2/PIB (Mw 35'000) 4:1:5 (Example 10) exhibited reversiblethermal transitions at the onset temperatures of 170° C. (16 J/g), 178°C. (17 J/g), 172° C. (6.2 J/g) and 169° C. (6.1 J/g) respectively. Thesecould be assigned to the ‘melting’ (deaggregation) and ‘crystallization’(aggregation) of the β-sheet aggregates. FIG. 18B providingThermogravimetric analysis revealed that all materials except PIB-NH2and Smactane™ stable up to temperatures of at least 250° C. Complete PIBdepolymerization occurred at temperatures above 340° C. for M2, D2 andtheir blends.

FIG. 19 Atomic force microscopy (AFM) height (left) and phase (right)images of M2 and M2/PIB (MW 35'000) 5:5 (Example 9) drop-cast fromtetrachlorethane solution onto SiO₂ substrates revealed the formation ofβ-sheet fibrils, while AFM images of pure PIB (MW 35'000) did not showsuch features.

FIGS. 20A, 20B, 20C, 20D, 20E and 20F provide Rheologicaltime-temperature superposition (TTS) master curves of M2/PIB (MW 75'000)5:5 (Example 8), M2/PIB (MW 35'000) 5:5 (Example 9), M2/D2/PIB (MW35'000) 4:1:5 (Example 10), in comparison to unmodified higher molecularweight PIB (MW 35'000, MW 75'000, MW 200'000) as well as Smactane™(hollow symbols in other graphs) as reference materials at T_(ref)=25°C. The blends gave rise to soft materials with a loss factor of tanδ>0.6 over almost the whole frequency range investigated.

FIG. 21A provides an Experimental setup for a random vibration modalanalysis test on a sandwich structure comprising a damping layer. Thefirst resonance frequency of the steel structure at 32.9 Hz and itsintensity decrease in the sandwich structure for PIB (MW 200′00), M2,for M2/D2/PIB (75 k) 5:5 (Example 8), M2/D2/PIB (35 k) 5:5 (Example 9),M2/D2/PIB (35 k) 4:1:5 (Example 10) as well as Smactane™. FIG. 21Bprovides Finite element simulations of the same sandwich configurationsfor damping layers based on the same materials as well as additionalcommercial damping materials. FIG. 21C provides experimental (circles)and calculated (squares) loss moduli G″ considered as the “figure ofmerit” for the vibration damping ability in constrained layers plottedrelative to the damping ratio with an exponential fit presented hereonly as a guide. These experimental G″ values were taken at 200 rad/sfrom a classical rheological frequency sweep test at 25° C.

FIGS. 22A, 22B, 22C, 22D, 22E, 22F, 22G and 22H provide Lissajous curvesobtained from oscillatory shear stress-strain test, performed on arheometer at −45° C. and 25 rad/s with a theoretical γ (% strain) valueimposed of 0.05%. FIG. 22A is for Smactane™ with γ=0.05%, FIG. 22B isfor PIB (Mw 200'000) with γ=0.045%, FIG. 22C is for M2 with γ=0.05%,FIG. 22D is for M2/PIB (MW 35'000) 5:5 (Example 9) with γ=0.068%, andFIG. 22E is for M2/D2/PIB (MW 35'000) 4:1:5 (Example 10) with γ=0.043%.The areas of the latter represent dissipated energies during one cycleof these tests. FIGS. 22F, 22G and 22H provide Plots of the logarithm ofthe dissipated energies relative to the logarithm of the strain appliedat −45° C. during an oscillatory shear stress-strain test, performed ona rheometer. Based on the mathematical equation: W_(d)=nG″ ε₀ ², fittingequations of PIB (Mw 200'000) (see FIG. 22F), M2/PIB (MW 35'000) 5:5(Example 9) (see FIG. 22G), and M2/D2/PIB (MW 35'000) 4:1:5 (Example 10)(see FIG. 22H) were used in order to correct the values of energiesdissipated for a γ (% strain) value of 0.05% for all the five materials.

Binary and ternary blends of compounds with “non-matching” oligopeptides(different oligopeptide length or amino acid sequence, includingnon-functionalised polymers) were found to give rise to novel“interpenetrating supramolecular networks”. Specifically, binary blends(compositions by weight) of M2/D1 1:4 (Example 7), M2/polyisobutylene MW75'000 (Example 8), M2/polyisobutylene MW 35'000 (Example 9), as well asthe ternary blend M2/D2/polyisobutylene (MW 35'000) 4:1:5 (Example 10)were obtained by dissolving mixtures of the compounds intetrachlorethane (TCE), stirring the solutions at room temperature for16 h, removing the solvent in vacuo, and drying the resulting materialsin high vacuum at 120° C. for 3 days. The annealing temperature waschosen such that it was below the melting transition of the tape-formingcomponent M2 but above the softening temperature of the secondnetwork-forming component (D1 or polyisobutylene). In the case ofExample 7, the gelation point of M2 and the entanglement point of D1were superimposed in their blends, resulting in materials with both agelation and an entanglement point within a similar frequency range (ina classical rheological frequency sweep at 25° C.). Thus the materialexhibited a large frequency region with pronounced liquid-like behaviour(that is, G″>G′) at room temperature, confined by two regions of solidelastomer-like (G′>G″) behaviour at higher and lower shear frequencies,as seen from a rheological time-temperature superposition (TTS) mastercurve at 25° C. (all shift factors log a_(T) and log b_(T). for the TTSmaster curves listed in Table 2). Whilst high molecular weightpolyisobutylene materials may exhibit such regions of “liquid-like”behaviour at temperatures just above their glass transition temperatureof T_(g)≈−65° C., Example 7 showed such behaviour over a largetemperature range and a broad frequency range at room temperature,exhibiting a large loss factor of up to tan δ=2.0 in this region, whichis unprecedented in related materials. Upon heating Example 7 above thematerials' glass-transition temperature of about T_(g)=−55° C., both itsstorage and loss moduli as well as viscosity first closely followedthose of D1. At about 65° C., where pure D1 is already in its liquidregime, the storage and loss moduli became similar to those of M2 andremained constant up to the melting transition at above 139° C. Thepronounced effect of the minority component M2, the two-stagetemperature transition, and the superimposed rheological properties ofthe pure components in the blend provide sufficient evidence for thepresence of two independent hydrogen-bonded superstructures, resultingin an “interpenetrating supramolecular network”. Whereas the highfrequency boundary of the “liquid-like” region can be assigned to thehydrogen-bonded network formed by D1, the low frequency boundary isassociated to a percolation network of the M2 β-sheet tapes.

The observed pronounced “liquid-like” behaviour (in a certaintemperature and mechanical frequency range), that is, the apparentmolecular level properties approaching liquid-like properties (such asflow) in a solid and macroscopically shape-persistent polymer materialsis the prerequisite for the self-healing properties of the materialsdisclosed here.

For the various blends of M2, D2, and polyisobutylenes (Examples 8-10),Solution-phase IR spectra in chlorinated solvents showed that allmixtures remained non-aggregated in solution. This enabled us to obtainhomogenous blends from solution, so as to obtain hydrogen-bondedaggregates from M2 dispersed in PIB as a matrix material (Examples 8-9)that can be cross-linked using the difunctional network-formingdifunctional D2 (Example 10). Solid-state infrared (IR) spectroscopy ofthe bulk materials revealed strongly aggregated and highly orderedantiparallel β-sheet structures. Thermogravimetric analysis (TGA) anddifferential scanning calorimetry (DSC) proved that the materials werethermally stable (against degradation) up to at least 250° C. Allmaterials exhibited sharp and reversible thermal transitions attemperatures of 169-178° C. Comparing the enthalpies of fusion of pureM2 and D2 (16-17 J/g) to those of the blends (6-6.5 J/g), we concludedthat 75-80% of the oligopeptide-modified components were aggregated intoβ-sheet tapes or fibrils in the blends. Moreover, temperature-dependentsolid-state IR spectroscopy on the materials proved that the observedtransition was associated to β-sheet deaggregation in all cases.Visualization of the nanoscopic morphologies of the obtained aggregatesby means of atomic force microscopy (AFM) imaging of continuous 1 μmthick films drop-cast from TCE solution onto SiO₂ substrates proved thatβ-sheet tape or fibril structures were present in those bulk materials.

In order to evaluate the mechanical properties of Examples 8-10, wetested their rheological properties in comparison to unmodified highermolecular weight polyisobutylenes (MW 35'000, 75'000, and 200,000).Compared to Smactane™, PIB (MW 200'000) showed lower storage (G′) andloss moduli (G″) over the whole range of frequencies (10⁻⁴-10⁶ rad/s) ortemperatures (−45° C.-105° C.) investigated, but a slightly higher andbroader peak of the loss factor tan δ (as a function of frequency),corresponding to its glass transition. Likewise, pure M2 and the blendsM2/PIB (Examples 8-9) showed significantly lower storage and loss moduliover a frequency range of 10⁻⁴-10³ rad/s, but significantly higher lossfactors of tan δ≈1 for a large frequency range. Examples 8 and 9, forinstance, gave rise to a soft rubber-like material with storage and lossmoduli G′ and G″ that were very similar to one another over almost thecomplete range of investigated frequencies, as determined from a TTSmaster curve at room temperature (all shift factors log a_(T) and logb_(T). for the TTS master curves listed in Table 2). As a result, theloss factor of Examples 8 and 9 peaked at tan δ=1.1 at a reducedfrequency of about a_(T)ω=5×10⁴ rad/s and never fell below tan δ=0.6 inthe reduced frequency range of a_(T)ω=10⁻³-10⁶ rad/s. The ternary blendM2/D2/PIB (MW 35'000) 4:1:5 (Example 10) possesses higher storage andloss moduli as well as similar loss factors compared to Smactane™, butwith even improved moduli and loss factor at low frequencies, due to theaddition of D2 which acts as a network forming crosslinker, resulting inan extra reinforcement of the materials. Moreover, the obtained mastercurves had a substantially different shape as compared to either pure M2or unmodified polyisobutylenes. Specifically, the rubbery plateau in thelow-frequency regime was absent, indicating that neither does M2 justserve as a filler, nor does the polyisobutylene matrix just act as adiluting “solvent”. One can therefore attribute the large temperatureand frequency range of high loss factor tan δ values to aninterpenetration of the PIB entanglement network and a percolationnetwork formed by the M2 β-sheet tapes. The resulting supramolecularnetworks exhibit an improved vibration damping performance wasattributed to improved energy dissipation by the high fraction ofpendant polymer chains incorporated into the network.

Such “interpenetrating supramolecular networks” as described hereprovide an alternative to traditional IPNs for the preparation ofhigh-performance vibration damping materials. In order to evaluate theperformance of Examples 8-10 with other damping materials, we testedtheir shear vibration damping characteristics of the in comparison tounmodified higher molecular weight polyisobutylenes (MW 35'000, 75'000,and 200,000), as well as Smactane™, a commercially available highperformance damping material with excellent damping propertiesspecifically at low temperatures. To this end, we employed a randomvibration modal analysis test on a sandwich structure representing atypical constrained damping layer application. The test structure wasdesigned specifically to investigate the structural damping performanceof the material in low frequency vibration (30-40 Hz), which is typicalfor the first vibration modes of many steel or aluminium panels used inautomotive or aerospace applications. The specimen with a free length of54 mm was fixed at one end while a weight of 4.6 g was clipped to theother one. The beam was excited using a pseudo random signal using avibration shaker driven through an open loop random vibrationcontroller. Accelerometers were used to monitor the base and tipaccelerations and reconstruct the frequency response function of thesystem around its first resonance peak. The modal damping ratio wasobtained by single-degree-of-freedom modal curve fitting of theresonance peak in the complex domain.

All results of the damping tests and derived damping ratios are listedin Table 3.

The first resonance frequency for steel, in this particular set up,occurred at 32 Hz and was slightly damped by steel itself and itsclamping on the base (0.4%). However, damping ratio significantlyincreased to 2.9% once Smactane™ was used as the damping layer in thesandwich structure. By comparison, while unmodified high molecularweight PIB (MW 200'000) exhibited a low damping ratio of 1.4%, thedamping ratios were 3.2% for pure M2, 2.6% for the binary blend M2/PIB(MW 75'000) (Example 8), 2.5% for the binary blend M2/PIB (MW 35'000)(Example 9), and 3.4% for the ternary blend M2/D2/PIB (MW 35'000)(Example 10). Examples 9 and 10 thus showed excellent damping ratios,even exceeding those of the commercially available high-performancedamping material Smactane™ and by far surpassing those of unmodified PIBthat is considered to possess good damping properties and is alreadyused in damping applications on a technological scale.

Moreover, we complemented our results with detailed finite element (FE)simulations of the sandwich beam vibration tests. We performed thesimulations also on other high-performance damping materials as areference, including Smactane™, Soundcoat™ Dyad 601 and 3M ISD™ 130.54using their rheology data as the input to the FE simulations. Inqualitative agreement with the experimental results, the finite elementsimulations resulted in damping ratios of Examples 8-10 above those ofthe reference materials.

The damping properties of the investigated materials at low temperatureswere obtained using the loss moduli G″ obtained from rheologicalfrequency sweep experiments at those temperatures and calculating thedissipated energies during one cycle of oscillatory stress-strain testfrom the area of the corresponding Lissajous curves (corrected for theimposed strain).

The calculated damping properties at low temperatures are listed inTable 4.

All of the investigated materials and, in particular, the ternary blendM2/D2/PIB (MW 35'0000) (Example 10) exhibit excellent damping propertiesat temperatures of −45° C. and below, down to their glass transitiontemperatures at about T_(g)≈−65° C., even exceeding those of Smactane™and in marked contrast to PIB (MW 200,000) that possesses lower lossmoduli over the whole range of frequencies (0.1-100 rad/s) tested at−45° C.

All commercial reference materials (except PIB) are composites withformulations highly optimized for damping performance. It is worthnoting that, as a consequence, our materials were light (with a densityof 0.92 g/cm³ compared to 1.18 g/cm³ for Smactane™) and did not requireany additional fillers or low molecular weight plasticizers. Moreover,as the “ideal” damping characteristics depend on the application, theversatility offered by the use of oligopeptide-modified polymers asadditives to commercial elastomers appears to provide an excellentpathway towards lightweight, low-creep, and high-performance constrainedlayers for vibro-acoustic damping.

Referring in particular to FIGS. 23A-25 (Examples 11-13):

FIGS. 23A, 23B, 23C and 23D show AFM phase images and differentialscanning calorimetry measurements that prove that oligopeptideaggregation competes with polymer phase segregation in blends of M2/S2and M3/S3 and leads to nanoscale phase segregation. As shown in FIG.23A, blends of M2 with unmodified polystyrene (9:1) as a referencematerial domains (white) with average sizes of hundreds of nanometers indiameter, only very few such polystyrene domains could be observed inblends M2/S2 (9:1) (Example 11) that were otherwise mostly homogeneous.As shown in FIG. 23B, in differential scanning calorimetry measurements,the glass transition of PS could be identified in both cases,demonstrating that even in the blend M2/S2 (9:1) (Example 11) nanoscalepolystyrene domains are present. As shown in FIG. 23C, blends of M3 withunmodified polystyrene (9:1) as a reference material forms polystyrenedomains (white) with average sizes of dozens of nanometers in diameter,and no such polystyrene domains were observed in blend M3/S3 (9:1)(Example 12). As shown in FIG. 23D, in this case, the polystyrene glasstransition was only observed in the differential scanning calorimetrymeasurements of the reference materials and not in the blend M3/S3 (9:1)(Example 12), proving that the polystyrene domain size was now so smallthat no cooperative properties such as a glass transition were observed.

FIGS. 24A and 24B show a) storage moduli G′ (FIG. 24A) and b) lossmoduli G″ (FIG. 24B) determined by rheological dynamic frequency sweepexperiments at 25° C. of M2, S2, the blend of M2 with unmodifiedpolystyrene (9:1) and the blend M2/S2 9:1 (Example 11); FIGS. 24C and24D show storage moduli G′ and loss moduli G″ determined by rheologicaldynamic frequency sweep experiments at 25° C. of M3, S3, the blend of M3with unmodified polystyrene (9:1) and the blend M3/S3 9:1 (Example 12).Notably, whereas the storage and loss moduli of the blends of either M2or M3 with unmodified polystyrene (9:1) increased slightly compared tothe pure M2 and M3, respectively, a strong increase in the storage andloss moduli was observed for the two blends M2/S2 9:1 (Example 11) andM3/S3 9:1 (Example 12). In particular, in the low frequency range, theaddition of 10 wt % of S3 in the blend M3/S3 9:1 (Example 12) issufficient to achieve a ten-fold increase in storage modulus G′. Thisproves that the polystyrene domains in these blends served to “gluetogether” the oligopeptide aggregates (tapes and nanofibrils), givingrise to a new percolation network of hard domains within the material.

FIG. 25 shows a rheological time-temperature superposition master curveat a reference temperature of 25° C. of the ternary blend M3/S3/PIB (MW35'000) 9:3:12 (Example 13). The storage moduli G′, loss moduli G″ andthe loss factor tan δ are reported for a large range of reducedfrequency (10⁶-10⁻⁴ rad s⁻¹), as well as temperature range (−45° C. to65° C.). The γ (% strain) values used for these experiments was 0.1% andwere selected to be in the linear regime of the investigated materials.The TTS master curve revealed that the tan δ was higher than 0.35 over alarge frequency range of 10⁻³-10⁻⁵ rad s⁻¹, and the G″ was higher than 1Mpa for frequencies higher than 10 rad s⁻¹. This combination of a largetemperature range of a large loss factor and a high storage modulus isideal for good, temperature-invariant damping properties. Accordingly,the material showed a very high damping ratio of =3.3% in constrainedlayer damping tests on a sandwich structure at a resonance frequency of32 Hz, as described in previous examples.

2. PREPARATIVE EXAMPLES 2.1 Instrumentation and Methods

NMR Spectroscopy was carried out on a Bruker Avance 300 spectrometeroperating at a frequency of 300.23 MHz for ¹H and 75.49 MHz for ¹³Cnuclei, or on a Bruker Avance 400 spectrometer operating at a frequencyof 400.23 MHz for ¹H and 100.63 MHz for ¹³C nuclei. Deuterated solventswere purchased from Cambridge Isotope Laboratories. The spectra werecalibrated to the respective residual proton peaks of the deuteratedsolvents (¹H NMR: 7.26 ppm CDCl₃, 6.0 ppm TCE-d₂, 5.32 ppm, DMSO-D₆,3.31 ppm CD₃OD; ¹³C NMR: 77.16 ppm CDCl₃, 49.00 ppm CD₃OD, 39.52 ppmDMSO-D₆).

Solution—Phase FTIR Spectra were recorded on a “Spectrum One” IRspectrometer from Perkin Elmer using a solution-phase cuvette with KBrwindows and a light path of 0.5 mm, or on a Jasco FT/IR-6300 Fourier,using a KBr window 32×3 mm (by Pike Technology). The materials weredissolved in either TCE or CHCl₃ (5 mg/mL). The TCE solutions werestirred for 1 h at 100° C. and left to slowly cool to room temperature.The CHCl₃ solutions were stirred at RT for 1-16 h.

Solid-State FTIR Spectra were recorded on a Bruker ALPHA FTIRspectrometer or on a JASCO FT/IR 6300 spectrometer using the Miracle ATRaccessory from PIKE, as well as on a Varian Fourier Transformspectrometer equipped with a Golden Gate diamond ATR with temperaturecontrol up to 200° C.

High Resolution Mass Spectra were recorded at the Mass SpectrometryService of EPFL on either an AXIMA Performance device from ShimadzuBiotech for MALDI-TOF and a Q-TOF Ultima from Waters for ESI-TOF, or atthe Mass Spectrometry Service of ETH Zurich on a Bruker Daltonics maXisfor HiRes-ESI-MS.

Thermogravimetric Analyses were performed on a TGA Q500 device from TAInstruments, loaded with samples of more than 2 mg. The measurementrange was 50-710° C. in a nitrogen atmosphere or 50-914° C. in air. Aheating rate of 10° C./min was applied in all cases.

Differential Scanning calorimetry was performed on a DSC Q1000 from TAInstruments in a nitrogen atmosphere, loaded with samples of more than 2mg. For the measurement of glass transition or melting temperatures,three measurements were performed in the range of −80 to 400° C. Boththe heating and cooling rates were 10° C./min. All data were collectedfrom the second heating cycle.

Combustion Elemental Analyses were carried out as service measurementsat EPFL using EA 1100 CHN Instrument or at the Institute of OrganicChemistry at ETH Zurich using a LECO CHN/900 instrument.

AFM Imaging was performed on a Nanoscope IIIa instrument. Samples wereprepared from stock solutions of the compounds in tetrachlorethane (TCE)at an initial concentration c=10⁻³ mol/L. The solutions were placed intosealed tubes, vigorously stirred (400 rpm) and heated in an oil bath to180° C. for 2 h, followed by stepwise cooling 160° C. (1 h), 140° C. (1h), 120° C. (1 h), and 100° C. (1 h) under continued stirring (100 rpm).Afterwards, the heating was switched off and the solution was allowed tocool to room temperature at the same stirring rate. The solutions werethen diluted to a concentration of c=1×10⁴ mol/L or c=5×10⁻⁵ mol/L andspin-coated onto SiO₂ substrates treated with ethanol and ultrapurewater (3000 rpm) or onto freshly cleaved HOPG (1800 rpm). The obtainedsamples were analyzed in tapping mode at room temperature in air, usingcantilevers with an average resonance frequency of 75 kHz and scan ratesof 0.5-1.5 Hz. The image resolution was 512×512 pixels.

Dynamic Shear Rheology Measurements were carried out on parallel platerheometers AR 2000, ARES LR2 or ARES from TA Instruments. Disc shapedsample specimen from all materials that were shape-persistent wereprepared on a Rittal table press. A force of 2 kN was applied for 30 minat 100° C., after which the specimen were cooled to 20° C. at a force of1.1 kN for 30 min. Depending on the samples quantity, aluminium platesof 15 mm or 25 mm diameter, as well as stainless steel plates of 25 mmdiameter were used. Discs of 12 mm diameter were prepared as well andplaced in the centre of the stainless-steel plates (25 mm diameter) witha centring tool. The gap between the plates was in the range of 0.4-2mm. In the case of non-adhesive samples, the plates were covered withemery paper to avoid wall slipping. Measurements were carried out attemperatures of −45° C. to 250° C. Once the desired temperature wasreached, the system was equilibrated for 2 min. Frequency sweeps rangingfrom 100 rad s⁻¹ to 0.01 rad s⁻¹ were carried out under controlledstrain. Depending on the sample composition and temperature, the strainamplitude ranged from 0.03% to 50%. The applied strain was defined suchthat the sample stayed in its linear viscoelastic domain during thecomplete frequency sweep.

Modal Damping Tests were carried out on a RMS 3000 vibration shakerusing an HP 35670A vibration controller and signal analyzer in open looppseudo random vibration analysis. The acceleration of the base and ofthe tip of the specimen were monitored using two Bruel&Kaer 4517accelerometer through a B&K Nexus 2692 amplifier. The frequency transferfunctions of the different specimens were measured in a frequency rangeof 10 to 110 Hz with a resolution of 0.125 Hz. The test specimenconsisted of a sandwich structure representing a constrained dampinglayer application. The base substrate was a steel plate with dimensionsof 60×6×0.5 mm onto which a damping layer with dimensions of 40×0.6×2.2mm was superimposed. A thin steel plate with dimensions of 40×0.6×0.2 mmwas used to constrain the top of the damping layer. The sandwich teststructure was then clamped with one end to the vibration table over alength of 6 mm, and a mass of 4.6 g was added to the free end of thebeam over a length of 5 mm. The added mass has been calculated such thatthe first bending mode of the sandwich beam is in the range of 30-40 Hz.In order to determine the modal damping ratio the first peak of thefrequency response function was first fitted using a complex polynomialfraction least square method integrated in the HP 35670A signalanalyzer. The modal damping ratio was then calculated from the real partX and the imaginary part ω of the first complex pole of the polynomialfraction using the definition ξ=λ/ω.

Finite Element Simulations of the sandwich beam vibration tests havebeen carried out to compare the damping performance of M2/PIB with otherhigh-performance damping materials in this particular application. Thechosen specimen geometry was the same as the sandwich beam specimen usedin the experimental modal damping tests (60×6×0.5 mm steel base platewith a 40×0.6×2.2 mm damping layer constrained on top by a 40×0.6×0.2 mmsteel sheet). The base plate was modelled as being clamped to the shakeron one side (imposed displacement, no rotation, over 6 mm) and attachedto two steel blocks of 5×5×15 mm. The steel plates were modelled aslinear elastic with a Young modulus of 210 GPa, Poisson ratio of 0.3 andmass density of 7,800 kg/m³. The damping layer materials were allconsidered incompressible and modelled using an Arruda-Boycehyperelastic potential (power exponent of 7) with complex shear modulitaken from rheology measurements or literature data (Table 3). The wholespecimen was modelled using 2640 3D hybrid quadratic hexahedric elements(14345 nodes) in Simulia Abaqus© 6.10 (FIG. 4e in the main paper) andsubjected to a steady state harmonic simulation in a frequency range of20-70 Hz with a resolution of 0.333 Hz to capture in detail the firstresonance peak at around 34 Hz. The simulated frequency responsefunctions (tip vs base acceleration in the direction of the Y axis) werethen processed to extract the modal damping ratio ξ (Table 3) using asingle degree of freedom complex curve fitting method.

2.2 Materials and General Synthesis Procedures Materials.

Reagents were purchased as reagent grade from commercial sources andused without further purification. Poly(isobutylene) 1 (Kerocom™ PIBA)was obtained from BASF and purified from non-functionalizedpoly(isobutylene) by column chromatography prior to use.Poly(isobutylene) diamine 4 containing about 10% of monofunctionalpoly(isobutylene) amine was obtained from BASF SE, Germany, and usedwithout further purification. THF, acetonitrile, toluene,dichloromethane and triethylamine were purchased as HPLC grade and driedusing a solvent purification system from Innovative technologies. Othersolvents were purchased as reagent grade and distilled once prior touse. Thin Layer Chromatography (TLC) Analyses were performed on TLCplates from Merck; UV-light (254 nm) or standard colouring reagents wereused for detection. Column Chromatography was conducted on Geduran®Silica gel Si 60 from Merck (40-60 μm).

Sample Preparation. For the preparation of films and solid samples ofeither single compounds or blends, the oligopeptide-polymer derivativesand/or PIB (MW 75,000) were dissolved in either TCE or CHCl₃, thesolutions were stirred at room temperature for 1-16 h, and concentratedin vacuo. The resulting materials were dried in HV at 120° C. for 3days.

General Procedure A: Peptide Coupling. The carboxylic acid derivativewas dissolved in THF. The amine (1 equiv) was added, as well asN-ethyldiisopropylamine (DIEA; 3 equiv) and(benzotriazol-1-yloxy)tripyrrolidinophosphonium hexafluorophosphate(PyBOP; 1.2 equiv). The solution was stirred for 3-16 h, and thereaction progress was monitored by TLC. The crude product was typicallypurified by precipitation into water (see General Procedures C or D).Specific purification or sample preparation procedures were performedbefore further characterization in some cases.

General Procedure B: Fmoc Deprotection. The Fmoc-protected aminederivatives were dissolved in CHCl₃. Then, a large excess of piperidine(≥15 equiv) was added, and the solution was stirred overnight. Thereaction progress was monitored by TLC. After completion of thereaction, the solvents were removed in vacuo. Unless otherwise noted,the crude product was purified by column chromatography.

General Procedure C: Precipitation of Compounds Soluble in THF. Aftercompletion of the reaction affording the desired compound, the reactionmixture was concentrated to half of its original volume. A large excessof aqueous 1 M HCl solution was added. The resulting precipitate wasfiltered off, re-dissolved in THF and precipitated again, following thesame procedure as described above. After three repetitiveprecipitations, the crude product was finally dissolved in CH₂Cl₂, CHCl₃or THF. The solution was dried over MgSO₄ and concentrated in vacuo at40° C.

General Procedure D: Precipitation of Compounds Insoluble in THF. Aftercompletion of the reaction affording the desired compound, the reactionmixture was diluted with a large excess of aqueous 1 M HCl solution. Theprecipitate was collected and re-dispersed in THF at 60° C. The productwas precipitated again using the same procedure as described above twomore times. The precipitate was finally re-dispersed in THF andconcentrated in vacuo at 40° C.

2.3 Synthesis Procedures and Analytical Data for 2-3, 5-6, M0-M5 andD0-D5

Synthesis of PIB₁₉-Ala₃-Fmoc 2. Following General Procedure A, PIB₁₉-NH₂1 (14.55 g, 12.03 mmol) andN-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (6.0 g,13.23 mmol) were dissolved in THF (250 mL). DIEA (6.18 mL, 36.08 mmol)and PyBOP (7.51 g, 14.43 mmol) were added. The reaction mixture wasstirred overnight. The product was precipitated following GeneralProcedure C. The final product (19.2 g, 97%) was obtained as a slightlyyellow wax. ¹H NMR (400 MHz, CDCl₃) δ=7.75 (d, J=7.5 Hz, 2H, aromaticH), 7.57 (d, J=7.5 Hz, 2H, aromatic H), 7.52 (m, 1H, NH), 7.38 (d, J=7.4Hz, 2H, aromatic H), 7.29 (t, J=7.4 Hz, 2H, aromatic H), 7.13 (m, 1H,NH), 6.76 (m, 1H, NH), 5.84 (m, 1H, NH), 4.73-4.29 (m, 5H, Fmoc-CO₂CH₂,3 CHCH₃), 4.20 (t, J=7.0 Hz, 2H, fluorenyl CH), 3.31-3.07 (m, 2H,CH₂NH), 1.81-0.53 (m, 178H, aliphatic H, 3 CHCH₃). MS (MALDI-TOF,DCTB/NaTFA 10:1): calcd for C₇₃H₁₂₆N₄O₅Na: (n=10[M+Na]+) 1161.9620;found: 1161.8517. R_(f): 0.45 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N₂)T_(g)=−66° C.

Synthesis of PIB₁₉-Ala₃-H 3. Following General Procedure B,PIB₁₉-Ala₃-Fmoc 2 (17.2 g, 10.46 mmol) was dissolved in CHCl₃ (250 mL).Piperidine (10.35 ml, 104.57 mmol) was added, and the reaction mixturewas stirred at room temperature overnight. The crude product waspurified by column chromatography (silica gel, gradientCH₂Cl₂→CH₂Cl₂/MeOH 5:1). The final product (10 g, 67%) was obtained as aslightly yellow wax. 1H NMR (400 MHz, CDCl₃) δ=7.83 (d, J=7.3 Hz, 1H,NH), 7.10 (m, 1H, NH), 6.46 (m, 1H, NH), 4.45 (m, 2H, CHCH₃), 3.51 (q,J=6.9 Hz, 1H, CHCH₃NH₂), 3.4-3.1 (m, 2H, CH₂NH), 1.82-0.65 (m, 178H,aliphatic H, 3 CHCH₃). MS (MALDI-TOF, DHB): calcd for C₅₈H₁₁₆N₄O₃Na:(n=10 [M+Na]⁺) 939.8940; found: 940.0981. R_(f): 0.15 (CH₂Cl₂/MeOH10:1). DSC (10° C./min, N₂) T_(g)=−68° C.

Synthesis of Fmoc-Ala₃-PIB₄₀-Ala₃-Fmoc 5. Following General Procedure A,N-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (4.36 g,9.61 mmol) and NH₂-PIB₄₀-NH₂ 4 (10.9 g, 4.28 mmol) were dissolved in THF(400 mL). DIEA (2.47 mL, 14.42 mmol) and PyBOP (5.50 g, 10.58 mmol) wereadded. After 16 h, the crude product was precipitated following GeneralProcedure D. The final product (14.18 g, 94%) was obtained as a whitesolid. ¹H NMR (400 MHz, C₂D₂Cl₄ at 110° C.) δ=7.81 (d, J=7.5 Hz, 4H,aromatic H), 7.63 (d, J=7.3 Hz, 4H, aromatic H), 7.50-7.41 (m, 5H,aromatic H), 7.4-7.3 (m, 4H, aromatic H), 7.21 (m, 3H, aromatic H), 6.52(m, 2H, NH), 6.34 (m, 2H, NH), 6.00 (m, 2H, NH), 5.16 (m, 2H, NH),4.63-4.08 (m, 12H, 6 CHCH₃, 2 Fmoc-CO₂CH₂, 2 fluorenyl CH), 3.37-2.94(m, 4H, 2CH₂NH), 1.94 (s, 4H, 2CH₂C(CH₃)₂Ph), 1.62-0.98 (m, 346H,aliphatic H, 6 CHCH₃), 0.95 (s, 12H, 2PhC(CH₃)2). MS (MALDI-TOF, DCTB):calcd for C₁₀₄H₁₅₈N₈O₁₀Na: (n+m=9 [M+Na]⁺) 1702.1993; found: 1702.3004.R_(f): 0.4 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N₂) T_(g)=−55° C.

Synthesis of H-Ala₃-PIB₄₀-Ala₃-H 6. Following General Procedure B,Fmoc-Ala₃-PIB₄₀-Ala₃-Fmoc 5 (11.00 g, 3.2 mmol) was dissolved in CHCl₃(200 mL). Piperidine (200 ml, 2.02 mol) was added, and the reactionmixture was stirred at room temperature overnight. The next day, thesolvent was evaporated in vacuo, and the mixture was washed three timeswith cold heptane. The crude product was then dispersed in DCM andconcentrated in vacuo at 40° C. Finally, the product (7.76 g, 82%) wasobtained as a white solid. ¹H NMR (400 MHz, CDCl₃ and TFA) δ=7.96 (m,2H, NH), 7.71 (m, 6H, NH2+TFA), 7.47 (m, 2H, NH), 7.38 (m, 1H, aromaticH), 7.21-7.06 (m, 3H, aromatic H), 6.88 (m, 2H, NH), 4.7-4.2 (m, 4H,CHCH₃), 3.41-2.79 (m, 6H, 2 CH₂NH, 1 CH₂NH₂), 1.84 (s, 4H,2CH₂C(CH₃)₂Ph), 1.77-0.87 (m, 346H, aliphatic H, 6CHCH₃), 0.80 (s, 12H,2 PhC(CH₃)₂). MS (MALDI-TOF, DHB): calcd for C₈₂H₁₅₅N₈O₆: (n+m=11[M+H]⁺) 1348.2065; found: 1348.5963. R_(f): 0.05 (CH₂Cl₂/MeOH 10:1). DSC(10° C./min, N₂) T_(g)=−65° C.

Synthesis of PIB₁₉-Ac M0. PIB₁₉—NH₂ 1 (3.85 g, 3.18 mmol) was dissolvedin THF (80 mL). Acetyl chloride (0.454 mL, 6.37 mmol) and pyridine(0.642 mL, 7.96 mmol) were added. The reaction mixture was stirredovernight. The crude product was precipitated following GeneralProcedure C. The final product (3.53 g, 88%) was obtained as a slightlyyellow viscous oil. ¹H NMR (400 MHz, CDCl₃) δ=5.36 (m, 1H, NH),3.47-3.08 (m, 2H, CH₂NH), 1.96 (s, 3H, C═OCH₃), 1.73-0.60 (m, 169H,aliphatic H). MS (MALDI-TOF, DCTB/NaTFA 10:1): calcd for C₅₁H₁₀₃NONa(n=10 [M+Na]⁺) 768.7932; found 768.9546. R_(f): 0.85 (CH₂Cl₂/MeOH 10:1).DSC (10° C./min, N2) T_(g)=−67° C.

Synthesis of PIB₁₉-Ala-Ac M1. Following General Procedure A, PIB₁₉-NH₂ 1(3.69 g, 3.05 mmol) and N-acetyl-L-alanine (0.4 g, 3.05 mmol) weredissolved in THF (100 mL). DIEA (1.57 mL, 9.15 mmol) and PyBOP (1.9 g,3.66 mmol) were added.

The reaction mixture was stirred overnight. The product was precipitatedfollowing General Procedure C. The final product (3.9 g, 97%) wasobtained as a slightly yellow viscous oil. ¹H NMR (400 MHz, CDCl₃ andTFA) δ=7.91 (d, J=7.6 Hz, 1H, NH), 6.87 (m, 1H, NH), 4.75-4.36 (m, 1H,CHCH₃), 3.30 (m, 2H, CH₂NH), 2.16 (s, 3H, C═OCH₃), 1.64-0.84 (m, 172H,aliphatic H, 1 CHCH₃). MS (MALDI-TOF, DCTB/NaTFA 10:1): calcd forC₅₄H₁₀₈N₂O₂Na: (n=10 [M+Na]⁺) 839.8303; found: 839.6927. R_(f): 0.55(CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N₂) T_(g)=−65° C.

Synthesis of PIB₁₉-Ala2-Ac M2. Following General Procedure A, PIB₁₉—NH₂1 (17.48 g, 14.46 mmol) and N-acetyl-L-alanyl-L-alanine (3.8 g, 18.79mmol) were dissolved in THF (250 mL). DIEA (4.95 mL, 28.91 mmol) andPyBOP (9.03 g, 17.35 mmol) were added. The reaction mixture was stirredovernight. The product was precipitated following General Procedure C.The final product (19.5 g, 97%) was obtained as a slightly yellow wax.¹H NMR (400 MHz, CDCl₃) δ=7.44 (m, 1H, NH), 7.00 (m, 1H, NH), 6.74 (m,1H, NH), 4.70 (m, 1H, CHCH₃), 4.59 (m, 1H, CHCH₃), 3.45-3.09 (m, 2H,CH₂NH), 2.05 (s, 3H, C═OCH₃), 1.62-0.64 (m, 175H, aliphatic H, 2 CHCH₃).MS (MALDI-TOF, CHCA/NaTFA 1:1): calcd for C₅₇H₁₁₃N₃O₃Na: (n=10 [M+Na]⁺)910.8674; found: 910.7816. R_(f): 0.4 (CH₂Cl₂/MeOH 10:1). DSC (10°C./min, N2) T_(g)=−68° C., T_(m)=170° C.

Synthesis of PIB₁₉-Ala₃-Ac M3. PIB₁₉-Ala₃-H 3 (2.99 g, 2.1 mmol) wasdissolved in THF (150 mL). Acetyl chloride (0.3 mL, 4.2 mmol) andpyridine (0.424 mL, 5.25 mmol) were added and the reaction mixture wasstirred overnight. The next day, the product was precipitated followingGeneral Procedure C. The final product (3.0 g, 98%) was obtained as aslightly yellow wax. ¹H NMR (400 MHz, CDCl₃ and TFA) δ=7.58 (m, 2H, NH),7.13 (m, 1H, NH), 6.90 (m, 1H, NH), 4.66-4.55 (m, 3H, CHCH₃), 3.26 (m,2H, CH₂NH), 2.11 (s, 3H, C═OCH₃), 1.70-0.57 (m, 178H, aliphatic H, 3CHCH₃). MS (MALDI-TOF, CHCA/NaTFA 1:1): calcd for C₆₀H₁₁₈N₄O₄Na: (n=10[M+Na]⁺) 981.9045; found: 981.8898. R_(f): 0.25 (CH2Cl2/MeOH 10:1). DSC(10° C./min, N2) T_(g)=−67° C.

Synthesis of PIB₁₉-Ala₄-Ac M4. Following General Procedure A,PIB₁₉-Ala₃-H 3 (2.0 g, 1.41 mmol) and N-acetyl-L-alanine (184.36 mg,1.41 mmol) were dissolved in THF (200 mL). DIEA (0.722 mL, 4.22 mmol)and PyBOP (877.98 mg, 1.69 mmol) were added. The reaction mixture wasstirred overnight. The product was precipitated following GeneralProcedure C. The final product (2.0 g, 90%) was obtained as a whiterubber. ¹H NMR (400 MHz, CDCl₃ and TFA) δ=7.83 (d, J=7.4 Hz, 1H, NH),7.60 (m, 1H, NH), 7.34 (d, J=6.1 Hz, 1H, NH), 7.27 (m, 1H, NH), 6.80 (m,1H, NH), 4.89-4.22 (m, 4H, CHCH₃), 3.45-3.09 (m, 2H, CH₂NH), 2.14 (s,3H, C═OCH₃), 1.65-0.74 (m, 181H, aliphatic H, 4 CHCH₃). MS (MALDI-TOF,CHCA/NaTFA 1:1): calcd for C₆₃H₁₂₃N₅O₅Na: (n=10 [M+Na]⁺) 1052.9416;found: 1052.8131. R_(f): 0.1 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N2)T_(g)=−66° C.

Synthesis of PIB₁₉-Ala₅-Ac M5. Following General Procedure A,PIB₁₉-Ala₃-H 3 (2.0 g, 1.41 mmol) and N-acetyl-L-alanyl-L-alanine(284.29 mg, 1.41 mmol) were dissolved in THF (200 mL). DIEA (0.722 mL,4.22 mmol) and PyBOP (877.98 mg, 1.69 mmol) were added. The reactionmixture was stirred overnight. The product was precipitated followingGeneral Procedure C. The final product (2.2 g, 95%) was obtained as awhite rubber. ¹H NMR (400 MHz, C₂D₂Cl₄ and TFA at 65° C.) δ=7.44 (d,J=6.9 Hz, 1H, NH), 7.16 (m, 3H, NH), 6.81 (m, 1H, NH), 6.56 (m, 1H, NH),4.67-4.42 (m, 5H, CHCH₃), 3.45-3.09 (m, 2H, CH₂NH), 2.17 (s, 3H,C═OCH₃), 1.84-0.68 (m, 184H, aliphatic H, 5 CHCH₃). MS (MALDI-TOF,CHCA): calcd for C₆₆H₁₂₈N₆O₆Na: (n=10 [M+Na]⁺) 1123.9788; found:1124.3684. R_(f): 0.05 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N2)T_(g)=−69° C.

Synthesis of Ac-PIB₄₀-Ac D0. NH₂-PIB₄₀-NH₂ 4 (5.00 g, 1.96 mmol) wasdissolved in THF (200 mL). Acetyl chloride (0.80 mL, 11.22 mmol) andpyridine (0.9 mL, 11.22 mmol) were added, and the solution was stirredovernight. The next day, the crude product was precipitated followingGeneral Procedure C. The final product (4.4 g, 85%) was obtained as acolourless viscous oil. ¹H NMR (400 MHz, CDCl₃) δ=7.37 (m, 1H, aromaticH), 7.23-7.06 (m, 3H, aromatic H), 5.44 (m, 2H, NH), 3.2-2.9 (m, 4H, 2CH₂NH), 1.98 (s, 6H, C═OCH₃), 1.84 (s, 4H, 2CH₂C(CH₃)₂Ph), 1.75-0.86 (m,328H, aliphatic H), 0.80 (s, 12H, 2 PhC(CH₃)₂). MS (MALDI-TOF, DCTB):calcd for C₇₆H₁₄₄N₂O₂Na: (n+m=13 [M+Na]⁺) 1140.1120; found: 1139.5068.R_(f): 0.7 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N2) T_(g)=−55° C.

Synthesis of Ac-Ala-PIB₄₀-Ala-Ac D1. Following General Procedure A,N-acetyl-L-alanine (810.0 mg, 6.17 mmol) and NH₂-PIB₄₀-NH₂ 4 (7.00 g,2.75 mmol) were dissolved in THF (200 mL). DIEA (1.58 mL, 9.26 mmol) andPyBOP (3.53 g, 6.79 mmol) were added. After 16 h, the crude product wasprecipitated following General Procedure C. The final product (5.12 g,66%) was obtained as a yellow glue. ¹H NMR (400 MHz, CDCl₃) δ=7.37 (m,1H, aromatic H), 7.23-7.06 (m, 3H, aromatic H), 6.17 (m, 4H, NH), 4.45(m, 2H, CHCH₃), 3.25-2.9 (m, 4H, 2 CH₂NH), 2.00 (s, 6H, C═OCH₃), 1.84(s, 4H, 2CH₂C(CH₃)₂Ph), 1.76-0.85 (m, 334H, aliphatic H, 2 CHCH₃), 0.80(s, 12H, 2 PhC(CH₃)₂). MS (MALDI-TOF, DCTB): calcd for C₈₆H₁₆₂N₄O₄Na:(n+m=14 [M+Na]⁺) 1338.2488; found: 1338.2838. R_(f): 0.5 (CH₂Cl₂/MeOH10:1). DSC (10° C./min, N₂) T_(g)=−54° C.

Synthesis of Ac-Ala₂-PIB₄₀-Ala₂-Ac D2. Following General Procedure A,N-acetyl-L-alanyl-L-alanine (700.0 mg, 3.46 mmol) and NH₂-PIB₄₀—NH₂ 4(3.93 g, 1.54 mmol) were dissolved in THF (300 mL). DIEA (0.89 mL, 5.19mmol) and PyBOP (1.98 g, 3.81 mmol) were added, and the reaction mixturewas stirred overnight. The next day, the crude product was precipitatedfollowing General Procedure D. The final product (4.0 g, 89%) wasobtained as an off-white solid. ¹H NMR (400 MHz, CDCl₃ and TFA)δ=7.64-7.3 (m, 6H, NH), 7.37 (m, 1H, aromatic H), 7.23-7.06 (m, 3H,aromatic H), 4.7-4.5 (m, 4H, CHCH₃), 3.25-2.9 (m, 4H, 2 CH₂NH), 2.07 (s,6H, C═OCH₃), 1.84 (s, 4H, 2CH₂C(CH₃)₂Ph), 1.77-0.86 (m, 300H, aliphaticH, 4 CHCH₃), 0.80 (s, 12H, 2 PhC(CH₃)₂). MS (MALDI-TOF, DCTB): calcd forC₉₂H₁₇₂N₆O₆Na: (n+m=14 [M+Na]⁺) 1480.3231; found: 1480.6152. R_(f): 0.4(CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N2) T_(g)=−57° C., T_(m)=178° C.

Synthesis of Ac-Ala₃-PIB₄₀-Ala₃-Ac D3. H-Ala₃-PIB₄₀-Ala₃-H 6 (1.50 g,0.50 mmol) was dissolved in THF (200 mL). Acetyl chloride (0.16 mL, 2.23mmol) and pyridine (0.18 mL, 2.23 mmol) were added, and the reactionmixture was stirred overnight. The next day, the crude product wasprecipitated General Procedure D. The final product (1.43 g, 92%) wasobtained as an off-white solid. ¹H NMR (400 MHz, CDCl₃ and TFA)δ=8.1−7.5 (m, 5H, NH), 7.38 (m, 2H, aromatic H, NH), 7.21-7.06 (m, 3H,aromatic H), 6.83 (m, 2H, NH), 4.7-4.3 (m, 6H, CHCH₃), 3.35-2.95 (m, 4H,2 CH₂NH), 2.12 (s, 6H, C═OCH₃), 1.84 (s, 4H, 2CH₂C(CH₃)₂Ph), 1.77-0.86(m, 346H, aliphatic H, 6 CHCH₃), 0.80 (s, 12H, 2 PhC(CH₃)₂). MS(ESI-TOF): calcd for C₇₄H₁₃₅N₈O₈Na: (n=8 [M+H+Na]²⁺) 643.5145; found:646.5165. R_(f): 0.3 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N2) T_(g)=−55°C.

Synthesis of Ac-Ala₄-PIB₄₀-Ala₄-Ac D4. Following General Procedure A,N-acetyl-L-alanine (146.0 mg, 1.11 mmol) and H-Ala₃-PIB₄₀-Ala₃-H 6 (1.50g, 0.50 mmol) were dissolved in THF (200 mL). DIEA (215.83 mL, 1.67mmol) and PyBOP (637.3 mg, 1.22 mmol) were added. After 16 h, the crudeproduct was precipitated following General Procedure D. The finalproduct (1.49 g, 91%) was obtained as an off-white solid. ¹H NMR (400MHz, CDCl₃ and TFA) δ=8.1-7.5 (m, 6H, NH), 7.38 (m, 2H, aromatic H, NH),7.21-7.06 (m, 4H, 3 aromatic H, NH), 6.83 (m, 2H, NH), 4.7-4.35 (m, 8H,CHCH₃), 3.35-2.95 (m, 4H, 2 CH₂NH), 2.12 (s, 6H, C═OCH₃), 1.84 (s, 4H,2CH₂C(CH₃)₂Ph), 1.77-0.86 (m, 352H, aliphatic H, 8 CHCH₃), 0.80 (s, 12H,2 PhC(CH₃)₂). MS (ESI-TOF): calcd for C₉₂H₁₇₀N₁₀O₁₀: (n+m=11 [M+2H]²⁺)787.6545; found: 787.1820. R_(f): 0.3 (CH₂Cl₂/MeOH 10:1). DSC (10°C./min, N2) T_(g)=−55° C.

Synthesis of Ac-Ala₅-PIB₄₀-Ala₅-Ac D5. Following General Procedure A,N-acetyl-L-alanyl-L-alanine (195.1 mg, 0.96 mmol) andH-Ala₃-PIB₄₀-Ala₃-H 6 (1.30 g, 0.43 mmol) were dissolved in THF (200mL). DIEA (0.25 mL, 1.45 mmol) and PyBOP (552.3 mg, 1.06 mmol) wereadded. After 16 h, the crude product was precipitated following GeneralProcedure D. The final product (1.30 g, 88%) was obtained as off-whitesolid. ¹H NMR (400 MHz, CDCl₃ and TFA) δ=8.1-7.5 (m, 10H, NH), 7.38 (m,2H, aromatic H, NH), 7.21-7.06 (m, 3H, aromatic H), 6.85 (m, 1H, NH),4.7-4.2 (m, 10H, CHCH₃), 3.35-2.95 (m, 4H, 2 CH₂NH), 2.12 (s, 6H,C═OCH₃), 1.84 (s, 4H, 2CH₂C(CH₃)₂Ph), 1.77-0.86 (m, 358H, aliphatic H,10 CHCH₃), 0.80 (s, 12H, 2PhC(CH₃)₂). MS (ESI-TOF): calcd forC₉₈H₁₇₈N₁₂O₁₂Na: (n+m=11 [M+Na]⁺) 1739.3613; found: 1740.5134. R_(f):0.25 (CH₂Cl₂/MeOH 10:1). DSC (10° C./min, N₂) T_(g)=−54° C.

Synthesis of PS₁₅-NH₂ 7. Styrene (12.19 g, 117.1 mmol) was freshlydistilled from CaH₂ prior to use. A rigorously dried 250 mL Schlenkflask was filled with dry cyclohexane (50 mL), and sec-butyl lithium(5.6 mL, 7.8 mmol, 1.4 M solution in cyclohexane) was added slowly via asyringe at 10° C. Under vigorous stirring, styrene was added via asyringe as fast as possible, leading to a yellow to orange color of thesolution. The cooling bath was removed and the mixture was stirred for 1h. Then, 2.5 excess of1-β-bromopropyl)-2,2,5,5-tetramethyl-1-aza-2,5-disilacyclopentane (5.47g, 19.5 mmol) dissolved in dry THF (20 mL) was added via a syringe,causing an immediate decoloration. The mixture was stirred for 2 h atroom temperature, then was concentrated in vacuo and taken up in THF(100 mL). Then, 1M HCl (35 mL) was added, and stirring was continuedovernight. The mixture was concentrated in vacuo, taken up in CH₂Cl₂,washed twice with 1M KOH and once with sat. NaCl solution. The combinedorganic phases were dried over MgSO₄, filtered, and concentrated invacuo. The crude product was purified by column chromatography (silicagel, gradient CH₂Cl₂→CH₂Cl₂/MeOH 20:1). The amine terminated polystyrene(9.2 g, 70%) was obtained as a white solid. ¹H NMR (400 MHz, CDCl₃)δ=7.25-6.3 (m, 75H, Ph-H), 2.48 (m, 2H, CH₂NH₂), 2.36-0.5 (m, 58H, 15CH₂CHPh, 3 CH₂, 1 CHCH₃, 2 CH₃). MS (ESI-TOF): calcd for C₁₂₇H₁₃₈N:(n=15 [M+H]⁺) 1678.0858; found: 1677.9187. R_(f): 0.35 (CH₂Cl₂/MeOH10:1).

Synthesis of PS₁₅-Ala₃-Fmoc 8. Following General Procedure A, PS₁₅-NH₂(5.04 g, 3.01 mmol) andN-(9-fluorenylmethyloxycarbonyl)-L-alanyl-L-alanyl-L-alanine (1.50 g,3.31 mmol) were dissolved in THF (250 mL). DIEA (1.54 mL, 9.02 mmol) andPyBOP (1.88 g, 3.61 mmol) were added. The reaction mixture was stirredovernight. The product was precipitated following General Procedure C.The final product (6.0 g, 95%) was obtained as a pinkish solid. ¹H NMR(400 MHz, CDCl₃) δ=7.76 (d, J=7.8 Hz, 2H, aromatic H), 7.55 (d, J=7.5Hz, 2H, aromatic H), 7.41 (t, J=7.4 Hz, 2H, aromatic H), 7.32 (t, J=7.3Hz, 2H, aromatic H), δ=7.3-6.3 (m, 75H, Ph-H), 6.23 (s, 1H, NH), 5.20(s, 1H, NH), 4.50-4.29 (m, 5H, Fmoc-CO₂CH₂, 3 CHCH₃), 4.17 (m, 2H,fluorenyl CH), 3.02 (m, 2H, CH₂NH), 2.5-0.5 (m, 67H, 15 CH₂CHPh, 5 CH₂,3 CHCH₃, C═OCH₃).

Synthesis of PS₁₅-Ala₃-H 9. Following General Procedure B,PS₁₅-Ala₃-Fmoc (5.4 g, 2.56 mmol) was dissolved in CHCl₃ (250 mL).Piperidine (5.06 ml, 51.11 mmol) was added, and the reaction mixture wasstirred at room temperature overnight. The crude product was purified bycolumn chromatography (silica gel, gradient CH₂Cl₂→CH₂Cl₂/MeOH 50:1).The final product (3.55 g, 75%) was obtained as yellowish powder. ¹H NMR(400 MHz, CDCl₃) δ=7.70 (m, 1H, NH), δ=7.3−6.3 (m, 75H, Ph-H), 5.90 (m,1H, NH), 4.31 (m, 2H, CHCH₃), 3.45 (m, 1H, CHCH₃NH₂), 3.02 (m, 2H,CH₂NH), 2.5-0.5 (m, 67H, 15 CH₂CHPh, 5 CH₂, 3 CHCH₃, C═OCH₃).

Synthesis of PS₁₅-Ala₂-Ac S2. Following General Procedure A, PS₁₅-NH₂(3.0 g, 1.79 mmol) and N-acetyl-1-alanyl-1-alanine (0.36 g, 1.79 mmol)were dissolved in THF (250 mL). DIEA (1.22 mL, 7.15 mmol) and PyBOP (1.4g, 2.68 mmol) were added. The reaction mixture was stirred overnight.The product was precipitated following General Procedure C. The finalproduct (2.2 g, 67%) was obtained as pinkish solid. ¹H NMR (400 MHz,CDCl₃) δ=7.3−6.3 (m, 75H, Ph-H), 6.10 (d, J=6.1 Hz, 1H, NH), 5.79 (s,1H, NH), 4.42 (m, 1H, CHCH₃), 4.29 (m, 1H, CHCH₃), 3.02 (m, 2H, CH₂NH),2.5-0.5 (m, 67H, 15 CH₂CHPh, 5 CH₂, 3 CHCH₃, C═OCH₃).

Synthesis of PS₁₅-Ala₃-Ac S3. PS₁₅-Ala₃-H (1.2 g, 0.63 mmol) wasdissolved in THF (50 mL). Acetyl chloride (136 mL, 1.9 mmol) andpyridine (0.205 mL, 2.54 mmol) were added and the reaction mixture wasstirred overnight. The next day, the product was precipitated followingGeneral Procedure C. The final product (1.2 g, 98%) was obtained aswhite solid powder. ¹H NMR (400 MHz, CDCl₃ and TFA) δ=7.3-6.3 (m, 75H,Ph-H), 5.0-4.6 (m, 3H, NH), 3.75 (m, 3H, CHCH₃), 3.03 (m, 2H, CH₂NH),2.5-0.5 (m, 70H, 15 CH₂CHPh, 5 CH₂, 4 CHCH₃, C═OCH₃).

Synthesis of PS₁₅-Ala₄-Ac S4. Following General Procedure A, PS₁₅-Ala₃-H(602.8 mg, 0.319 mmol) and N-acetyl-1-alanine (46.0 mg, 0.351 mmol) weredissolved in THF (30 mL). DIEA (0.17 mL, 0.957 mmol) and PyBOP (602.9mg, 1.69 mmol) were added. The reaction mixture was stirred overnight.The product was precipitated following General Procedure C. The finalproduct (0.6 g, 95%) was obtained as a white solid powder. ¹H NMR (400MHz, CDCl₃ and TFA) δ=7.3-6.3 (m, 75H, Ph-H), 5.5-4.6 (m, 4H, CHCH₃),3.03 (m, 2H, CH₂NH), 2.5-0.5 (m, 73H, 15 CH₂CHPh, 5 CH₂, 5 CHCH₃,C═OCH₃).

TABLE 1 Representative rheological data of M0-M5, D0-D2,polyisobutylenes of different molecular weights, as well as the blendsM2/D2 99:1 (Example 1), M2/D2 95:5 (example 2), M2/D2 9:1 (Example 3),M2/D2 7:3 (Example 4), M2/D2 5:5 (Example 5), and M2/D2 1:9 (Example 6);storage moduli G′, loss moduli G″, loss factors tan δ, and shearviscosities |η*|. Materials G′/Pa G″/Pa tan δ |η*|Pa · s) M0 0.22 139623 139 M1 249 1250 5.02 1270 M2 42800 14300 0.334 45200 M3 251000 429000.171 254000 M4 510000 40200 0.079 511000 M5 633000 63900 0.101 636000D0 25 2690 109 2690 D1 12300 70700 5.75 71800 D2 2000000 117000 0.0592010000 PIB-NH₂ (MW 1200) 0.19 32 168 32 H₂N-PIB-NH₂ (MW 1200) 4.99 23446.9 235 PIB (MW 35′000) 13400 24600 1.84 28000 PIB (MW 75′000) 6450042800 0.664 77500 PIB (MW 200′000) 231000 25270 0.109 232000 PIB (MW425′000) 188100 11840 0.063 188400 M2/D2 99:1 123000 21800 0.177 125000M2/D2 95:5 564000 47310 0.084 565000 M2/D2 9:1 495000 29610 0.060 496000M2/D2 7:3 798000 52780 0.066 800000 M2/D2 5:5 1760000 124000 0.0701760000 M2/D2 1:9 1500000 93880 0.063 1500000

TABLE 2 Shift factors log a_(T) and log b_(T) and activation energiesE_(a) obtained from the Arrhenius equation for the rheologicalmeasurements used in the time-temperature superposition master curves ofpolyisobutylenes of different molecular weights, pure M2 and D1, as wellas the blends M2/D1 1:4 (Example 7), M2/PIB (MW 75′000) 5:5 (Example 8),M2/PIB (MW 35′000) 5:5 (Example 9), and M2/D2/PIB(MW 35′000) 4:1:5(Example 10). Materials T/° C. log a_(T) log b_(T) E_(a)/J mol⁻¹ PIB(200k) −45 4.474 0.119 1055.8 −10 1.677 0.028 25 0 0 105 −2.199 −0.111PIB (75k) −45 3.889 −0.17 986.91 −10 1.684 0.021 25 0 0 65 −1.317−0.0048 105 −2.271 −0.099 PIB (35k) −45 4.492 0.133 1080.16 −10 1.6890.04 25 0 0 65 −1.302 −0.025 105 −2.356 −0.166 M2 −45 4.317 −0.0891040.6 −10 1.65 −0.066 25 0 0 105 −2.255 0.015 D1 −25 3.901 −0.0331735.8 −10 2.991 0.137 25 0 0 −25 3.93 −0.024 M2/D1 1:4 −10 2.998 0.1121585.7 25 0 0 65 −1.918 0.115 105 −3.928 −0.119 −25 3.901 −0.033 M2/PIB(75k) 5:5 −45 4.305 0.033 1078.1 −10 1.701 0.035 25 0 0 65 −1.282 −0.024105 −2.573 −0.216 M2/PIB (35k) 5:5 −45 4.075 −0.145 1111.9 −10 1.478−0.065 25 0 0 65 −1.901 −0.364 105 −2.881 −0.464 M2/D2/PIB (35k) 4:1:5−45 4.4 0.037 1136.3 −10 1.6 0.057 25 0 0 65 −1.841 −0.186 105 −2.677−0.265 Smactane −45 4.342 −0.059 1194.3 −10 1.746 0.05 25 0 0 65 −1.8−0.035 105 −3.208 −0.045

TABLE 3 Storage moduli G′ and loss moduli G″ used as input in the finiteelement simulations, and the resonance frequencies ω and modal dampingratios ξ resulting from these simulations for the reference materialsSmactane ™, PIB (MW 200′000), Soundcoat ™ Dyad 601, and 3M ISD 130, aswell as M2 and the blends M2/PIB (MW 75′000) 5:5 (Example 8), M2/PIB (MW35′000) 5:5 (Example 9), and M2/D2/PIB(MW 35′000) 4:1:5 (Example 10).Materials G′/MPa G″/MPa ω/Hz ξ Smactane ™ 1.01 0.47 36.1 1.8% PIB (200k)0.317 0.148 35 0.70% Soundcoat Dyad 601² 1.26 0.756 36.5 2.6% 3M ISD130¹ 0.094 0.036 34.6 0.20% M2 0.424 0.524 35.3 2.3% M2/PIB (75k) 5:50.509 0.442 35.396 1.89% M2/PIB (35k) 5:5 0.377 0.371 35.1 1.7%M2/D2/PIB (35k) 4:1:5 1.48 0.83 36.8 2.8%

TABLE 4 Calculation of the low-temperature damping properties. Lossmodulus G″ and loss factor tan δ at 200 rad/s (≈32 Hz) from arheological frequency sweep at room temperature, first resonancefrequency ω in the ‘forced vibration tests’ on sandwich structures,modal damping ratio ζ at room temperature, and calculated dissipatedenergy W_(d) at −45° C. and 25 rad/s for steel, Smactane ™, PIB (MW200′000), M2, M2/PIB (MW 35′000) 5:5 (Example 9), and M2/D2/PIB(MW35′000) 4:1:5 (Example 10). G″/MPa tan δ ω/Hz ζ W_(d)/J m⁻³ Steel — —32.2 0.4% — Smactane ™ 0.47 0.45 34.9 2.9% 5.98 PIB (MW 200′000) 0.150.45 32.4 1.1% 3.95 M2 0.54 1.3 33.4 3.2% 13.29 M2/PIB (35k) 5:5 0.371.0 33.4 2.5% 11.09 M2/D2/PIB (35k) 4:1:5 0.83 0.55 35.6 3.4% 18.95

TABLE 5 Shift factors log a_(T) and log b_(T) for the rheologicalmeasurements used in the time-temperature superposition master curves ofblends M3/S3/PIB (MW 35′000) 9:3:12 (Example 13). Materials T/° C. loga_(T) log b_(T) M3/S3/PIB (35k) −45 4.220 −0.220 −25 2.856 −0.115 −101.600 −0.170 25 0 0 65 −2.454 −0.033

1. A polymer blend, comprising: a first oligopeptide-terminal polymercomponent selected from the group consisting of: a hydrophobic, flexiblepolymer having a glass transition temperature below 20° C. and only onemonodisperse oligopeptide end group, the monodisperse oligopeptide endgroup having 1 to 5 amino acid repeating units; and a hydrophobic,flexible polymer having a glass transition temperature below 20° C. andtwo monodisperse oligopeptide end groups; and at least one additionalpolymer component selected from the group consisting of: a hydrophobic,flexible polymer that is different from said first oligopeptide-terminalpolymer component and that has a glass transition temperature below 20°C. and only one monodisperse oligopeptide end group, the monodisperseoligopeptide end group having 1 to 5 amino acid repeating units; and ahydrophobic, flexible polymer that is different from said firstoligopeptide-terminal polymer component and that has a glass transitiontemperature below 20° C. and two monodisperse oligopeptide end groups.2. The polymer blend according to claim 1, wherein at least one of saidfirst oligopeptide-terminal polymer component and said at least oneadditional polymer component comprises repeating units selected from thegroup consisting of isobutylene, butadiene, siloxane, acrylate, andfluoropolymer units.
 3. The polymer blend according to claim 1, whereinat least one of said first oligopeptide-terminal polymer component andsaid at least one additional polymer component comprises one or more ofisobutylene, isoprene or styrene units.
 4. The polymer blend accordingto claim 1, wherein said first oligopeptide-terminal polymer componentand said at least one additional polymer component include ahydrophobic, flexible isobutylene polymer having a glass transitiontemperature below 20° C. and only one monodisperse oligopeptide endgroup, the monodisperse oligopeptide end group having 1 to 5 amino acidrepeating units, blended with a hydrophobic, flexible styrene polymerhaving a glass transition temperature below 20° C. and only onemonodisperse oligopeptide end group, the monodisperse oligopeptide endgroup having 1 to 5 amino acid repeating units.
 5. The polymer blendaccording to claim 1, wherein an oligopeptide moiety of said firstoligopeptide-terminal polymer component comprises L-alanine units. 6.The polymer blend according to claim 1, wherein a polymer segment of atleast one of said first oligopeptide-terminal polymer component and saidat least one additional polymer component is selected from the groupconsisting of: polyisobutylene, poly(isobutylene-co-isoprene),polyisoprene, polybutadiene, polysiloxane, polyacrylate,poly(ethylene-co-butylene), hydrogenated poly(isoprene), hydrogenatedpoly(butadiene), and a fluoropolymer.
 7. The polymer blend according toclaim 6, wherein said fluoropolymer ispoly(tetrafluoroethylene-co-ethylene).
 8. The polymer blend according toclaim 1, wherein said first oligopeptide-terminal polymer component andsaid at least one additional polymer component include a hydrophobic,flexible polyisobutylene polymer having one or two monodisperseoligopeptide end groups blended with a hydrophobic, flexible polystyrenepolymer having one or two monodisperse oligopeptide end groups.
 9. Thepolymer blend according to claim 1, in the form of a shape-persistentthermoplastic elastomer.
 10. The polymer blend according to claim 1,wherein each of said oligopeptide end groups of said firstoligopeptide-terminal polymer component and said at least one additionalpolymer component is the same.
 11. The polymer blend according to claim10, wherein each of said oligopeptide end groups of said firstoligopeptide-terminal polymer component and said at least one additionalpolymer component has the same 2 amino acid repeating units beyond itsterminal amide group.
 12. The polymer blend according to claim 1,comprising interpenetrating supramolecular polymer networks in which twoor more specific supramolecular interactions result in the formation oftwo or more independent, interpenetrating supramolecular networks withdifferent transition temperatures, that is, deaggregation temperatures.13. A vibration damping material comprising the polymer blend accordingto claim
 12. 14. The vibration damping material according to claim 13,being a composite material including one or more of the following: aplasticizer; and a reinforcing filler comprising carbon fibre, carbonblack, or silica particles.
 15. The vibration damping material accordingto claim 13, in a form adapted to reduce vibration within a vehicle, theform being a pad or other layer which can be interposed between membersof the vehicle subject to vibration.
 16. A vehicle which includes thevibration damping material according to claim
 13. 17. The vehicleaccording to claim 16, which is a motor vehicle or an aerospace vehicle.18. A method of vibration damping which involves use of the polymerblend according to claim 1 upon or within a structure or a vehicle. 19.A polymer blend, comprising: at least one hydrophobic, flexible polymerhaving a glass transition temperature below 20° C. and only onemonodisperse oligopeptide end group, the monodisperse oligopeptide endgroup having 1 to 5 amino acid repeating units; and at least onehydrophobic, flexible polymer having a glass transition temperaturebelow 20° C. and two monodisperse oligopeptide end groups.
 20. Thepolymer blend according to claim 19, wherein said at least onehydrophobic, flexible polymer having only one monodisperse oligopeptideend group and said at least one hydrophobic, flexible polymer having twomonodisperse oligopeptide end groups comprise the same type of polymersegment.